surface modification of molds and acessories for the glass industry

2014

Doctor of Philosophy Thesis in Mechanical Engineering, Production Technology branch, supervised by Professor Altino de Jesus

Roque Loureiro and Professor Albano Augusto Cavaleiro Rodrigues de Carvalho and submitted to the Mechanical Engineering

Department of the Faculty of Sciences and Technology of the University of Coimbra.

SURFACE MODIFICATION OF MOLDS AND ACESSORIES FOR THE GLASS INDUSTRY

Filipe Daniel Fernandes

IMAGEM

2014

Doctor of Philosophy Thesis in Mechanical Engineering, Production Technology branch, supervised by Professor Altino de Jesus

Roque Loureiro and Professor Albano Augusto Cavaleiro Rodrigues de Carvalho and submitted to the Mechanical Engineering

Department of the Faculty of Sciences and Technology of the University of Coimbra.

SURFACE MODIFICATION OF MOLDS AND ACESSORIES FOR THE GLASS INDUSTRY

Filipe Daniel Fernandes

IMAGEM

Bolsa de Doutoramento (SFRH/BD/68740/2010)

Acknowledgments

i

Acknowledgments

Many people have contributed and encouraged me along all the years of this research, I thank

all of them, and I hope I have gathered a bit of their best.

Above all, I want to express my sincere gratitude to my supervisors: Albano Cavaleiro and

Altino Loureiro, for their scientific guidance, commitment and availability during the

execution of this thesis. Their bright mind and wisdom brought to me many insights and

knowledge that was the key success for this thesis. I’m also extremely indebted to them for all

the opportunities I had of involvement with various enterprises and scientific research centers

as well as of visiting several wonderful places in the world in the framework of conferences

and collaborations.

I have had the pleasure of working with and learning from a huge number of other professors

and researchers in different countries. I would like to specially thank professors Amilcar

Ramalho, Tomas Polcar, Josep Guilemany and Jerzy Morgiel for all the collaboration,

scientific discussion and open access to their laboratories.

CEMUC in the person of its chairman, Professor Valdemar Fernandes, is greatly

acknowledged by the possibility of carrying out most of this work within this research center.

Thanks are also due to INTERMOLDE company for the support and collaboration in the

initial stage of the thesis. Thanks are also extended to TEandM Enterprise.

It is essential to emphasize the crucial role of IPN (Instituto Pedro Nunes) on my work and

results, by making available laboratorial equipment for the development and characterization

of the samples I produced.

Special thanks to my friends and research colleagues from ECAT, Surface Engineering and

Nanomaterials and Micromanufacturing groups of CEMUC. I have spent marvelous moments

in their companionship. These thanks are extended to all the young and senior members from

Acknowledgments

ii

the different international groups (Barcelona, Prague, Southampton and Guangxi (China)) I

visited, who welcomed me and helped me during my training periods there.

To my close friends thank you for all the great moments.

I would like to acknowledge the funding provided by the FCT (Fundação para a Ciência e

Tecnologia), through SFRH/BD/68740/2010 fellowship and the projects AUTOMATIN

“5380”, PLUNGETEC “13545” and PTDC/EME-TME/122116/2010.

Finally, I would like to thank my parents, grandparents and sister who permanently gave me

their support during this hard but reward time. Without their patience and good advices, I

would not have been able to take advantage of the opportunity I was given to improve my

education. This thesis is dedicated to them. Last but not the least, I thank you Cláudia for your

daily personal support and encouragement in all good and bad times.

Abstract

iii

Abstract

Coatings are frequently applied on molds and accessories for the glass industry in

order to restrict surface degradation such as oxidation, corrosion, abrasion and wear of the

structural material, thereby decreasing the maintenance costs and increasing the lifetime and

performance of the components. However, in order to obtain accurate lifetime expectancies

and performance of the coatings it is necessary to have a complete reliable understanding of

their properties.

This thesis is on the improvement of the surface properties and integrity of molds, in

order to increase their durability, through the application of different types of coatings. Two

methodologies were followed to reach such demands: (i) optimization of coatings currently

used in molds surface protection (Ni-based alloys deposited by Plasma Transferred Arc -

PTA); (ii) synthesis and characterization of new coatings with improved functionalities,

deposited by emergent deposition processes such as APS - Atmospheric Plasma Spraying

(effect of nanostructured ZrO2 additions on Ni-based alloy coatings) and DCRMS - Direct

Current Reactive Magnetron Sputtering (influence of V additions on the properties of

TiSi(V)N thin films).

The dilution of the substrate in PTA process was shown to strongly influence the

structure and, consequently, the hardness, the oxidation resistance and the tribological

behavior of the coatings; with increasing dilution, a detrimental effect on these properties was

observed due to the incorporation of base material. However, in relation to the tribological

behavior, a beneficial effect at high temperature was demonstrated due to the fast formation of

oxide layers which protect the coating surface against wear. The post-weld heat treatment

performed at coatings reduced the hardness of the partially melted and heat affected zones

without affecting the coatings hardness; whereas the coatings hardness and wear resistance

was improved with annealing treatment. Thus, the best performing coating could only be

achieved by a proper selection of the deposition conditions, in order to get the best

compromise between mechanical properties, high temperature oxidation behavior and wear

resistance of the coatings.

The impact promoted by nanostructured ZrO2 additions on the microstructure of a Ni-

based alloy depended on the way how the coatings were deposited by APS process, using: (i)

nanostructured ZrO2 and Ni-alloy powders previously mixed by mechanical alloying or (ii)

Abstract

iv

the same powders supplied separately. A homogeneous and compact microstructure with

small zirconia particles evenly distributed in the matrix was achieved in the first case, while a

porous microstructure, full of semi-melted Ni powders with large particles of ZrO2 entrapped

in their boundaries, suggesting a brittle behavior, was deposited in the second. In both cases

the hardness and wear behavior of ZrO2 rich coatings were improved in relation to the Ni-

based alloy. The coatings deposited from mechanically alloyed powders revealed to be much

more tribologically performing due to their compact structure and even distribution of

zirconia particles. All the APS coatings showed higher hardness values than the Ni-based

coatings deposited by PTA; however, their micro-abrasion resistance was worse, due to the

lack of cohesion between the powders.

The analysis by XRD of the structure of V rich TiSi(V)N coatings, deposited by

DCRMS, revealed that V incorporation in the TiSiN system shifted the peaks to higher

angles, indicating the formation of a substitutional solid solution based on TiN phase, where

Ti atoms are replaced by the smaller V ones. On the other hand, by similar reason, XRD of

TiSiN films revealed that a nanocomposite structure consisting of TiN grains enrobed by a Si-

N matrix was not formed. In fact, with Si addition a shift of the diffraction peaks of TiN phase

to higher angles was observed which, in combination with similar compressive residual

stresses, also supported the formation of a substitutional solid solution. V additions showed to

successfully improve the hardness and tribological behavior of TiSi(V)N films, as a result of

the substitutional solid solution formation. On the other hand, the formation of the V2O5 phase

during the sliding contact acts as a lubricious tribo-film, protecting the coating against wear.

The addition of Si or V strongly influenced in opposite directions the oxidation resistance of

the coatings. Si incorporations in TiN increased significantly the oxidation resistance of the

films, whereas the opposite occurred in V containing coatings. In this latter case, the rapid V

ions out-diffusion through the oxide scale inhibited the formation of a continuous protective

silicon oxide layer, which is responsible for the excellent oxidation behavior of TiSiN films.

V rich coatings showed lower oxidation resistance than PTA-deposited Ni-based coatings, but

superior hardness values and better tribological behavior was found in relation to both PTA

and APS deposited coatings.

Keywords: Glass Industry, Molds surface protection, Ni-based thick coatings, TiSi(V)N thin

films, self-lubricant coatings.

Resumo

v * PTA – Do inglês – Plasma Transferred Arc

** APS – Do inglês – Atmospheric Plasma Spraying

*** PVD – Do inglês – Physical Vapor Deposition

Resumo

Frequentemente são aplicados diferentes tipos de revestimentos em moldes e

acessórios para a indústria do vidro, de forma a atenuar a degradação das suas superfícies,

devida às condições extremas de oxidação, corrosão, abrasão e desgaste a que estão sujeitas.

Deste modo, é possível, diminuir os custos de manutenção e aumentar o tempo de vida e a

eficácia destes componentes. No entanto, de forma a maximizar o desempenho dos

revestimentos, é necessário compreender a inter-relação entre as suas caraterísticas e

propriedades funcionais.

Esta tese é dedicada à melhoria das propriedades da superfície dos moldes, de forma a

aumentar a sua durabilidade e desempenho, através da aplicação de diferentes tipos de

revestimentos. No sentido de enfrentar este desafio, foram seguidas duas metodologias

alternativas: (i) otimização de revestimentos já utilizados na proteção da superfície de moldes

(ligas à base de Ni depositadas por PTA* (plasma de arco transferido); (ii) síntese e

caracterização de novos revestimentos mais resistente e com melhores propriedades,

depositados por processos de deposição emergentes no mercado, tal como são os casos do

APS** - projeção em plasma atmosférico (efeito da adição de ZrO2 nanoestruturada numa

liga à base de níquel) e do PVD*** - deposição física em fase de vapor (influência da adição

de V nas propriedades de revestimentos cerâmicos do sistema TiSi(V)N).

A diluição do substrato, promovida pelo processo PTA, mostrou influenciar a estrutura

e, consequentemente, diminuir a dureza, a resistência à oxidação e a resistência ao desgaste à

temperatura ambiente dos revestimentos à base de Ni. No entanto, teve um efeito benéfico no

comportamento tribológico a temperaturas elevadas, permitindo a rápida formação de uma

camada de óxido, que protege a superfície contra o desgaste. O tratamento térmico efectuado

aos revestimentos após deposição reduziu a dureza das zonas termicamente afetada pelo calor

e parcialmente fundida. Além disso, o tratamento de envelhecimento realizado aos

revestimentos permitiu aumentar a sua dureza e a sua resistência ao desgaste. Assim, o

revestimento com melhor desempenho só pode ser obtido selecionando criteriosamente as

condições de deposição que originam o melhor compromisso entre propriedades mecânicas,

comportamento à oxidação a alta temperatura e resistência ao desgaste.

Resumo

vi * PTA – Do inglês – Plasma Transferred Arc



A adição de zircónia nano-estruturada à liga de níquel depositada por APS tem um

impacto diferente na microestrutura dos revestimentos, consoante são usados pós de Ni e

ZrO2 previamente misturados por síntese mecânica ou os mesmos pós alimentados

separadamente. Quando depositados com os pós preparados por síntese mecânica, os

revestimentos exibiram uma estrutura homogénea e compacta com pequenas partículas de

zircónia uniformemente distribuídas na matriz de Ni. Por outro lado, quando depositados

separadamente foi obtida uma estrutura muito porosa com pós de Ni semi-fundidos e com

partículas de ZrO2 inseridas na sua interface, sugerindo um comportamento frágil.

Independente do procedimento de deposição, em ambos os casos a dureza e a resistência ao

desgaste dos revestimentos foi melhorada com a adição de ZrO2. Contudo, os revestimentos

depositados usando pós produzidos por síntese mecânica revelaram possuir melhor

performance, devido à sua estrutura compacta e distribuição uniforme da ZrO2. Todos os

revestimentos depositados por APS revelaram possuir valores de dureza superiores aos

revestimentos depositados por PTA; no entanto, a sua resistência à abrasão mostrou ser

bastante inferior devido à falta de coesão entre os pós.

As análises de difracção de raios-X realizadas aos filmes TiSi(V)N ricos em V

revelaram que os picos de difracção se deslocam para ângulos superiores com o aumento do

teor em V, sugerindo a formação de uma solução sólida substitucional, com o V de menor

raio atómico a substituir átomos de Ti na estrutura de base do TiN. Por outro lado, a análise

realizada aos revestimentos de TiSiN revelou que a estrutura nanocompósita (grãos de TiN

envolvidos numa matriz de Si-N) não foi alcançada. Com efeito, a adição de Si ao sistema

TiN também deslocou os picos de difracção para ângulos superiores o que, considerando o

mesmo nível de tensões residuais medidos em todos os revestimentos, indicou a formação de

uma solução sólida substitucional onde os átomos mais pequenos de Si substituíram os

átomos de Ti na rede do TiN. A adição de V ao sistema TiSiN permitiu melhorar as

propriedades mecânicas e tribológicas devido à formação da solução sólida substitucional. A

melhoria das propriedades tribológicas está relacionada com a formação da fase V2O5, no

topo da superfície do filme, que age como lubrificante sólido e, portanto, protege o

revestimento do desgaste. A resistência à oxidação de revestimentos ricos em Si e V mostrou

ser influenciada em sentidos contrários por estes elementos. Assim, a adição de Si ao sistema

TiN aumentou significativamente a resistência à oxidação dos filmes enquanto a adição de V

ao sistema TiSiN promoveu uma diminuição abruta na resistência à oxidação. Esta

diminuição da resistência à oxidação foi atribuída à rápida difusão dos iões de V através da

Resumo

vii * PTA – Do inglês – Plasma Transferred Arc



camada de óxido formada, que inibiu a formação de uma camada contínua e protectora de

óxido de silício, contrariamente ao que acontece nos revestimentos TiSiN. Os filmes ricos em

V apresentaram menor resistência à oxidação do que os revestimentos espessos depositados

por PTA; no entanto, a sua dureza e resistência ao desgaste foram bastante superiores à dos

revestimentos espessos depositados por PTA ou APS.

Palavras-chave: Indústria vidreira, Protecção da superfície de moldes, Revestimentos espessos

de Níquel, Revestimentos finos de TiSi(V)N, Revestimentos auto-lubrificantes.

Index

ix

Index

Acknowledgments ....................................................................................................................... i

Abstract ...................................................................................................................................... iii

Resumo ....................................................................................................................................... v

Index .......................................................................................................................................... ix

List of Figures ............................................................................................................................ xi

List of Tables ........................................................................................................................... xiii

Nomenclature............................................................................................................................ xv

1. Introduction ......................................................................................................................... 1

List of papers that are the basis of this thesis ............................................................................. 7

2. Historical perspective .......................................................................................................... 9

3. Effect of the arc current variation on the properties of Ni-based coatings deposited by

PTA process .............................................................................................................................. 21

3.1. Introduction ................................................................................................................ 21

3.2. Microstructure and hardness of the coatings ............................................................. 22

3.3. Oxidation resistance and diffusion mechanisms ........................................................ 26

3.4. Tribological behavior ................................................................................................. 28

3.4.1. Micro-scale abrasion tests ....................................................................................... 28

3.4.2. Pin-on-disc tests ...................................................................................................... 30

4. Influence of nanostructured ZrO2 additions on the wear resistance of Ni-based alloy

coatings deposited by APS ....................................................................................................... 33

4.1. Introduction ................................................................................................................ 33

4.2. Microstructure and hardness of the coatings ............................................................. 34

4.3. Tribological behavior ................................................................................................. 35

5. Study of the influence of V additions on the properties of TiSi(V)N films deposited by

DC reactive magnetron sputtering ............................................................................................ 39

Index

x

5.1. Introduction ............................................................................................................... 39

5.2. Microstructure, residual stresses and mechanical properties .................................... 40

5.3. Oxidation resistance and diffusion mechanisms ....................................................... 44

5.4. Tribological behavior ................................................................................................ 47

6. Conclusions and Future Research ..................................................................................... 51

7. References ......................................................................................................................... 59

Annex A ………………………….…………………………………...………………..…… 71

Annex B …………………………………………………………………………………….. 87

Annex C …………………………………………………………………………………….. 97

Annex D …………………………………………………………………………………… 113

Annex E ………………………………………………………………………………….… 123

Annex F ………………………………………………………………………………….… 145

Annex G …………………………………………………………………………………… 159

Annex H …………………………………………………………………………………… 171

Annex I ………………………………………………………………………………..…… 187

List of Figures

xi

List of Figures

Figure 2.1 - Components used in glass containers production. ................................................ 10

Figure 3.1 - Dilution induced by PTA process using different weldments: specimen 1 - 100 A,

specimen 2 and 3 - 128A, specimen 4 – 140 A. ....................................................................... 23

Figure 3.2 - Optical micrographs of: a) specimen 1, b) specimen 2, c) specimen 4, d) interface

of specimen 1. ........................................................................................................................... 24

Figure 3.3 - Hardness profile across the interface of PTA specimens. .................................... 25

Figure 3.4 - Hardness profile across the interface after PWHT in specimen 4. ....................... 26

Figure 3.5 - Isothermal oxidation curves at 800 and 900 0C of the: a) coating deposited using

100 A arc current, b) coating deposited using 128 A arc current. ............................................ 27

Figure 3.6 - Specific wear rate of Ni-based coatings in as-deposited and annealed conditions

from Annex A and C, evaluated in a micro-scale abrasion equipment. ................................... 29

Figure 3.7 - Variation of the wear rate of the as deposited PTA coatings and the gray cast iron

with testing temperature. .......................................................................................................... 31

Figure 4.1 - SEM morphology of the: a) Ni-based alloy coating, b) nanostructured ZrO2

coating, c) and d) Ni+20% Dual and Ni+40% Dual, respectively, e) and f) Ni+20% MA and

Ni+40% MA, respectively. ....................................................................................................... 35

Figure 4.2 - Specific wear rate of coatings calculated from the reciprocating tribological tests.

.................................................................................................................................................. 37

Figure 4.3 - Evolution of the wear behavior as a function of the parameter “normal load ×

sliding distance”. ...................................................................................................................... 38

Figure 4.4 - Specific wear rate of coatings, calculated from the slope of a straight line fitted to

the points plotted in Figure 4.3. ................................................................................................ 38

Figure 5.1 - XRD patterns of TiSiN coatings with Si content ranging from 0 to 12.6 at.% and

S1-y coatings with V additions. ................................................................................................ 41

Figure 5.2 - a) Effect of Si content on the hardness and Young’s modulus for the TiN system.

b) Effect of V content on the hardness and Young’s modulus forcoating S1-0. ...................... 42

Figure 5.3 - a) Adhesion critical loads values for the dissimilar coatings. Coatings were

scratched as the normal force was increased linearly from 5 to 70N. Lc2 critical load for TiN

coating was > 70 N. b), c) First coating cracking and first chiping observed on the scratch

track of coating S2-4, respectively. .......................................................................................... 44

List of Figures

xii

Figure 5.4 - a) Thermogravimetric isothermal analysis of coatings exposed at different

temperatures. Cross section morphology of coating: a) S2-0 oxidized at 900 ºC during 1 h, b)

S2-8 oxidized at 600 ºC during 30 min. ................................................................................... 45

Figure 5.5 - Wear rate of the coatings tested against Al2O3 and HSS balls. ............................ 48

Figure 5.6 - Evolution of the wear behavior as a function of the parameter “normal load ×

sliding distance”. ...................................................................................................................... 49

Figure 5.7 - Specific wear rate of coatings, calculated from the slope of a straight line fitted to

the points plotted in Figure 5.6. ............................................................................................... 49

Figure 6.1 - Hardness of coatings studied in the aim of this thesis, deposited by different

deposition techniques. .............................................................................................................. 53

Figure 6.2 - Comparison of the oxidation weight gain of Ni-based coatings deposited by PTA

and TiSi(V)N ones deposited by DC reactive magnetron sputtering. ...................................... 54

Figure 6.3- Specific wear rate of coatings calculated from the micro-scale abrasion tests. .... 56

Figure 6.4 - Specific wear rate of coatings calculated from the pin-on-disc tests. .................. 57

Figure 6.5 - Specific wear rate of coatings calculated from the pin-on-disc tests performed at

high temperature. ..................................................................................................................... 57

List of Tables

xiii

List of Tables

Table 2.1 - Comparison of the characteristics of several hardfacing and thermal spraying

processes [9, 55-60]. ................................................................................................................. 15

Table 2.2 - Typical sequential tasks used in molds production. ............................................... 17

Table 5.1 - Coatings deposited and their nomenclature. .......................................................... 40

Nomenclature

xv

Nomenclature

APS - Atmospheric plasma spraying

BM - Base material

DC - Direct current

DCRMS - Direct current reactive magnetron sputtering

FS - Flame spraying

FZ - Fusion zone

GTAW - Gas tungsten arc welding

HAZ - Heat affected zone

HT - High temperature

HVOF - High velocity oxy-fuel

MA - Mechanical alloying

PMZ - Partial melted zone

PTA - Plasma transferred arc

PVD - Physical vapor deposition

PWHT - Post-weld heat treatment

SEM-EDS - Scanning electron microscopy with energy dispersive spectroscopy

STEM - Scanning transmission electron microscope

TGA - Thermal gravimetric analysis

WDS - Wavelength-dispersive spectroscopy

XRD - X-ray Diffraction

http://www.oerlikon.com/ecomaXL/files/metco_Plasma_Solutions_EN3.pdf&download=1

http://en.wikipedia.org/wiki/Gas_tungsten_arc_welding

http://onlinelibrary.wiley.com/doi/10.1111/j.1749-6632.1986.tb34539.x/abstract

http://serc.carleton.edu/research_education/geochemsheets/wds.html

http://web.pdx.edu/~pmoeck/phy381/Topic5a-XRD.pdf

Chapter I – Introduction

1 Filipe Daniel Fernandes

Chapter I

1. Introduction

The necessity for higher performance and increased efficiency of components working

in extreme harsh conditions has been met with the use of more advanced materials. Coatings

have been developed over the last decades as a potential solution for these applications due to

their superior properties when compared to bulk materials, allowing extending the

performance and lifetime of components. This was the case of the glass industry, where the

components are actually coated with hard and wear resistant materials in order to resist to the

direct contact with the melted glass which promotes very severe abrasion, corrosion and

fatigue at high temperature. The current European glass container market is approximately 30-

40 billions of containers. Due to globalization, there has been a growing demand for glass

containers production at reduced costs, with increasing shape complexity, putting an

increasing pressure on glass producers who are seeking molds which allow increasing

production rates with higher quality. In these very harsh conditions, further degradation of the

mold surfaces occurs due to more intense thermal cycles leading to higher mechanical strains

and plastic deformation with the consequent fatigue, oxidation and wear problems.

Furthermore, other common problems such as worn-out geometries or surface and base

material cracks, often lead to the breakdown of the coating after few working cycles. The

combined effect of the harsh conditions with a faulty design or a defective material,

associated with bad working practices, induces the premature failure of the components.

Chapter I - Introduction


Besides the direct costs of the molds, there is a direct impact on the production due to the

increasing machine down times, higher process instability, poor product quality and higher

maintenance costs. Therefore, any optimized solution to achieve the best performance of

molds should be focused on the proper selection of the base material, the coating and its

deposition process. Unfortunately, the lack of credible solutions have limited the application

of coated components in glass industry, hiding the great relevance and impact they could have

for the industrial development in this area. Thus, this thesis is devoted to the improvement of

surface properties and integrity of molds, in order to improve their durability, through either

the optimization of conventional coatings currently applied on mold surfaces or the

development of new coating systems deposited by emergent technologies. Both cases show

potential to over-support the in service conditions of the molds, allowing improving their

ratability, lifetime and performance.

Presently, the molds for glass industry are manufactured either in low cost materials,

such as cast iron and low carbon steel, offering a good relative performance at high

temperature, or in copper-nickel alloys (bronze), due to their excellent thermal conductivity

combined with reasonable hardness and wear resistance, which allow an increasing

production rate of glass containers. In some cases, specific zones of the molds are coated,

mainly with Fe, Co or Ni-based alloys, with the aim of increasing their service life, owing to

their excellent oxidation and mechanical performance under high temperature conditions.

Several hardfacing and thermal spraying processes have been used to coat the molds: i)

Plasma Transferred Arc (PTA), and Gas Tungsten Arc Welding (GTAW) and ii) Flame

Spraying (FS) and High Velocity oxy Fuel (HVOF), respectively. Hardfacing processes have

been preferentially used in the protection of molds surface, due to the strong metallurgical

bond promoted between the coating and substrate, condition required to avoid the catastrophic

failure in service. In thermal spraying processes only a mechanical bonding is established.

Among the hardfacing processes referred above, PTA has been the most versatile and

used, due to its ability to deposit very thick coatings with low porosity and high production

rates. However, the final quality of coatings is largely influenced by the process parameters;

their incorrect selection leads frequently to problems such as the formation of porosity or

cracks in both the coating and the substrate, lack of adhesion of the coating to the substrate

and changes in the chemical composition of the coating metal alloy, compromising the

performance in service. Hence, a reliable understanding of the influence of the deposition

parameters on the coatings properties is required, in order to optimize the lifetime

expectancies and performance of the coatings. Presently, the tendency is to increase the



dilution of the substrate to avoid adhesion problems and, therefore, not to compromise the

molds life, even if the mechanical properties, oxidation resistance and wear performance are

deteriorated. This will be the first issue in discussion in this thesis: how the increase in

dilution of the base material (cast iron), induced by the change in arc current, influences the

properties of Ni-based coatings deposited by the PTA process.

An alternative route to improve the quality and performance of molds is the use of

new, thick or thin coatings deposited by emerging technologies that can lead to significant

improvements on the mechanical properties. Atmospheric Plasma Spraying (APS) and

Physical Vapor Deposition (PVD) are good examples of these deposition processes which

allow to achieve either metallic, ceramic or composite coatings. Such technologies and

materials have shown already to be potential candidates for improving the performance of

mechanical components and parts used in other industries where high oxidation, corrosion

and wear resistances are required. When compared to traditional deposition processes (PTA,

GTAW and FS), their main disadvantage is related to their high cost of investment and

production and, in some cases, lower coatings adhesion. Therefore, any development on this

field for molds protection should be focused on solutions that withstand the severe conditions

to which molds are submitted, extending their service life sufficiently, in order to justify their

higher cost. This is the second part of this thesis: new solutions based on thick and thin

coatings, deposited by APS and PVD processes, respectively. In the first case, the influence of

nanostructured ZrO2 additions to Ni-based alloy coatings deposited by APS process was

studied. The idea of this study was to take advantage of the high hardness and oxidation

resistance of ZrO2 to improve the oxidation resistance and the mechanical and wear properties

of Ni-based coatings. In relation to other thermal spraying processes, such as HVOF, APS

was selected due to its high potential to deposit materials with high melting temperature

(which is the case of ZrO2), although lower particles velocities are achieved; furthermore, in

relation to the PTA process, APS normally gives rise to coatings with much higher hardness.

Apart from the higher cost, the main disadvantage of APS coatings is the need of their

subsequent fusion when applied in parts for the glass industry, in order to obtain compact and

more performing coatings; the application of this operation is not covered in this thesis.

On the other hand, PVD technique has been extensively used to develop new coating

materials, due to its versatility. This process allows "sweep" the chemical composition of a

coating system composed by two or more chemical elements, by changing the deposition

parameters, producing small quantities of material that can be characterized in order to

optimize the best formulation. Moreover, in comparison to thick coatings, PVD allows the



deposition of films with much better properties, particularly the mechanical properties, for

which it is common to have improvements of two to four times. However, despite the

excellent properties of the films and the versatility of the process, its application in

components for the glass industry has been postponed, since the much smaller thickness, in

the micrometer order, has been seen unattractive to the industrial eyes when compared to

thick coatings. However, the experience gained in the application of such PVD films in other

fields of the metalworking industries (for example in machining, stamping, etc), where the

life-time of the coated parts has been easily increased by one to two orders of magnitude,

allows predicting the success of these films in the protection of molds surfaces for the glass

industry. Over the last decades a large number of films (ceramic and composite) have been

developed for high temperature applications, among which the most studied systems were

nitrides, carbides and oxides of Ti, Cr or Al, whose properties (mechanical, oxidation,

corrosion, wear, etc) have been successfully improved by the addition of other chemical

elements. One of the biggest advantages of thin coatings over thick ones is the possibility of

using ceramics as protection material. In fact, ceramic materials cannot be used in the surface

protection of molds for glass parts, due to their low thermal conductivity, which impede the

efficient cooling of the base material, decreasing the productivity of the process. This

difficulty can be overcome if these materials are deposited as thin coatings.

In this thesis, the addition of a specific element (vanadium) to ceramic films (TiSiN),

deposited by DC reactive magnetron sputtering, was studied. TiSiN is one of the best PVD

deposited coatings for high temperature applications and, with V incorporation, it is expected

to improve the mechanical properties and, at the same time, make them self-lubricant at high

temperatures. Currently, in the glass containers production, the walls of the molds are

frequently wet with liquid lubricant to avoid the adhesion of the melted glass to the molds

surface facilitating the removal of glass containers. Frequently, these lubricants (graphite

based lubricants) contain small amounts of sulfur which combined to oxygen forms sulfuric

acid that erodes the mold surfaces creating holes and, therefore, defects that lead to the

premature failure of the mold. Although V additions have been reported to improve the

mechanical properties and the tribological performance of some binary and ternary coating

systems, at high temperature V rapidly diffuses to the surface leading to the loss of its

lubricious effect after a critical short time. Hence, TiSiN system was selected due to its very

high oxidation resistance, similar to the ternary systems CrAlN and TiAlN, where V

incorporation was already studied, but with much better mechanical properties. Moreover, if

deposited with a nanocomposite structure (TiN grains embedded in a Si3N4 matrix) by



changing the thickness of the antidiffusion barrier Si3N4 matrix, the control of V diffusion can

be achieved.

Thus, the objectives of this research can be summarized as follows:

i) To analyze the effect of the arc current variation on the substrate dilution and, therefore, on

the chemical composition, microstructure, oxidation and tribological performance of Ni-based

coatings deposited by PTA process over gray cast iron. The influence of an annealing

treatment on the morphology, mechanical properties, and wear performance of coatings will

be also discussed.

ii) To study the influence of the addition of nanostructured ZrO2 to Ni-based coatings,

deposited by APS process, on their microstructure, mechanical properties and tribological

performance. The properties of the composite coatings will be compared with those of the

unmodified Ni and nanostructured ZrO2 coatings.

iii) To study the influence of V additions on the microstructure, mechanical properties,

oxidation and tribological behavior of TiSi(V)N thin films, containing different Si and V

contents, deposited by DC reactive magnetron sputtering. The properties of the V rich

coatings will be compared with TiSiN, TiN and TiVN coatings deposited as reference; special

attention will be focused on the oxidation mechanisms occurring during high temperature

exposure.

Taking into account these objectives, besides this introductory part, this thesis is

organized in six chapters. In chapter II, a historical perspective is presented about the main

problems addressed in molds for glass containers production, as well as on the main materials

and protection processes currently and possible use on their fabrication. In chapter III, IV and

V, the main achievements resulting from the experimental work on the deposition and

characterization of the coatings will be analyzed in relation to the three deposition techniques

above presented. Finally, in chapter VI the outputs of this thesis, a comparative synthesis and

discussion of all the results and future research topics will be presented.

The structure of the thesis is based on a short presentation and synthesis of the

research work performed in the scope of each deposition technique, which is further

supported by the papers already published and/or under submission by the author. The

chapters dedicated to the experimental studies are as follows:



Chapter III is dedicated to establish the correlations between the process parameters of

Ni-based coatings deposited by PTA and the chemical composition, morphology, structure,

mechanical properties (hardness), oxidation resistance, mechanisms of oxidation, structural

and morphology evolution after annealing and tribological behavior (evaluated by micro-scale

abrasion and pin-on-disc-tests) of the coatings. This chapter is supported by the papers

presented in annexes A to E.

Chapter IV is focused on the study of the influence of nanostructured ZrO2 additions

on the wear resistance of Ni-based alloy coatings deposited by APS. This analysis includes

the characterization of structure, morphology, mechanical properties and tribological behavior

(evaluated in a reciprocating sliding pin-on-disk equipment and micro-scale abrasion) of the

coatings. This chapter is supported by the paper in Annex F.

Chapter V is concentrated on the analysis of TiSiN films deposited by DC reactive

magnetron sputtering with different Si and V contents. The influence of V on the chemical

composition, structure, mechanical properties, residual stresses, oxidation resistance and

mechanisms of oxidation, thermal stability and tribological performance of the coatings is

addressed. This chapter is supported by the papers in Annexes G to I.

List of papers that are the basis of this thesis



Annex A - F. Fernandes, B. Lopes, A. Cavaleiro, A. Ramalho, A. Loureiro, Effect of arc

current on microstructure and wear characteristics of a Ni-based coating deposited by PTA

on gray cast iron, Surface and Coatings Technology, 205 (2011) 4094-4106.

Annex B - F. Fernandes, A. Cavaleiro, A. Loureiro, Oxidation behavior of Ni-based coatings

deposited by PTA on grey cast iron, Surface and Coatings Technology, 207 (2012) 196-203.

Annex C - F. Fernandes, A. Ramalho, A. Loureiro, A. Cavaleiro, Wear resistance of a nickel-

based coating deposited by PTA on grey cast iron, International Journal of Surface Science

and Engineering, 6 (2012) 201-213.

Annex D - F. Fernandes, A. Ramalho, A. Loureiro, A. Cavaleiro, Mapping the micro-

abrasion resistance of a Ni-based coating deposited by PTA on gray cast iron, Wear, 292-293

(2012) 151-158.

Annex E - F. Fernandes, T. Polcar, A. Loureiro, A. Cavaleiro, Room and high temperature

tribological behavior of Ni-based coatings deposited by PTA on gray cast iron, (2014), under

review, “Tribology International”.

Annex F - F. Fernandes, A. Ramalho, A. Loureiro, J.M. Guilemany, M. Torrell, A. Cavaleiro,

Influence of nanostructured ZrO2 additions on the wear resistance of Ni-based alloy coatings

deposited by APS process, Wear, 303 (2013) 591-601.

Annex G - F. Fernandes, A. Loureiro, T. Polcar, A. Cavaleiro, The effect of increasing V

content on the structure, mechanical properties and oxidation resistance of Ti–Si–V–N films

deposited by DC reactive magnetron sputtering, Applied Surface Science, vol. 289, (2014)

114-123.

Annex H - F. Fernandes, J. Morgiel, T. Polcar, A. Cavaleiro, Oxidation and diffusion

processes during annealing of TiSi(V)N films, (2014), under review, “Thin Solid Films”.

http://www.journals.elsevier.com/surface-and-coatings-technology/




Annex I - F. Fernandes, T. Polcar, A. Cavaleiro, Tribological properties of self-lubricating

TiSiVN coatings at room temperature, (2014), accept for publication, “Surface and Coatings

Technology”, DOI: 10.1016/j.surfcoat.2014.10.016

Reprints of the papers were made with the written consent of the Publisher and can be found

in annex.



Chapter II - Historical Perspective


Chapter II In the current chapter, the base and coating materials, processes of deposition currently used in molds protection

and the main causes for failure problems are reviewed. Moreover, an overview of possible solutions, regarding

processes and materials, for improving the efficiency of molds is presented.

2. Historical perspective

Glass was discovered by the Phoenicians around the year of 3000 BC. The history tells

that glass was discovered when these explorers lit fires near the beaches and observed the

changes occurring to the sand under the action of the intense heat of the fire. Time passed and

Romans began to master the technique of glass production being able to produce some

rudimentary bottles, intended for the storage of a large variety of liquids. Over the following

centuries, glass containers progressed significantly from the design point of view, having been

regarded as a luxury item in the middle age. At that time, only few wealthiest families used

glass containers to the storage of perfumes and other liquids. The industrialization of glass

containers started in England in the year 1600 when the coal was introduced as a solid fuel.

The glass containers were originally produced manually, but due to the increasing market

demand associated to technological development, in the XIX century glass production started

to be mechanically automated. Glass containers were seen as an ecologic and versatile

product, giving rise to their rapid expansion through the market. Since then, to achieve the

required higher production rates to satisfy the market needs, the use of more advanced

materials was always the preferred solution. During the 20th century, materials science and



technology knowledge was introduced in the industry and, from the second half of the

century, researchers and engineers poured a part of the materials development efforts on the

modification of materials surface through coatings application. Their superior properties when

compared to bulk materials, allowed extending the performance and lifetime of components

working in severe conditions. Since then, application of coatings has been one of the preferred

ways to modify the surface of materials permitting to extend their use to new potential

applications.

In glass industry, molds and accessories are in direct contact with the melted glass

(temperature about 1050 ºC) being submitted to very severe abrasion, corrosion, wear and

fatigue at high temperature. Molds, mouthpieces and plungers used for glass production are

shown in Figure 2.1. These mechanical components started to be produced in gray cast iron

and low carbon steel due to their relatively good high-temperature performance, allowing

keeping a fairly acceptable thermal conductivity, at an extremely low cost [1, 2]. In fact, the

functionality of these materials was not only to support the severe environment high-

temperature conditions but also to provide an efficient heat transfer which allows to rapidly

cool down the melted glass in order to either decrease the time of containers production or

obtain products without glass distorting under its own weight [2, 3]. Therefore, later on,

copper-alloys (bronze-aluminum) have been introduced as mold material due to their superior

thermal conductivity combined with good properties of hardness and wear resistance [4].

These materials allowed increasing production rates, despite of their much higher cost when

compared to the previous ferrous alloys. Independently of the mold material, some cooling

channels and fins had been performed in some components in order to help the heat

extraction.

Figure 2.1 - Components used in glass containers production.

1



The higher oxidation and wear resistance required at high temperatures (HT) for

achieving higher production rates, in order to satisfy market demands, led to the application of

new high performing bulk materials, such as Ni cast alloys, that allowed extend service life of

components [5, 6]. However, despite of their excellent performance at high temperature, their

price of acquisition/production was extremely high which have hindered their widespread

application in glass industry. Nevertheless, their attractive properties at high temperatures

pushed to overcome the cost barrier by their application under the form of thick coatings on

components fabricated in cheaper base materials. Currently, the molds are made in a low

performing low cost material (such as the mentioned above, i.e. cast iron, low carbon steel

and copper alloys) but coated in specific zones with hard and more resistant HT alloys,

allowing to support the harsh service conditions [2, 3, 7, 8]. The base material provides the

necessary mechanical strength to resist to the overall applied loads and the thermal fatigue,

whereas the coatings allow extending these properties to an upper end of their performance

capabilities while protecting them against wear, corrosion, oxidation, abrasion, etc [9, 10].

Even if the material withstands high temperature conditions without coating, the life period of

the component can be enhanced. In conclusion, the advantages of coatings application can be

summarized as follows [9, 11, 12]:

Cost of the coating is significantly lower than that of a component integrally fabricated

in HT alloys.

Unique structures and microstructures can be achieved in coatings which are not

possible in bulk materials.

Very high flexibility concerning the chemical composition of the deposited material

allowing the best selection and consequent optimization for a specific

solicitation/application.

With coatings, surface properties can be improved keeping the required mechanical

properties in the bulk of the structural component.

A wide variety of coating materials, such as cobalt, iron and nickel alloys, has been

successfully used to extend the working life of molds [2, 13-15]. The selection of the

appropriate coating composition depends on the environment to which the coating is exposed

and the substrate on which it is applied. The complexity of interactions between environment,

coating and substrate makes the design and selection of coatings very hard and, in general, a



compromise should be achieved between the mechanical strength, the corrosion/oxidation

resistance and the adhesion. Among the above referred coatings, nickel-based superalloys

have gained an important role for the protection of surfaces of components for the glass

industry, due to their excellent performance under conditions of abrasion, corrosion, wear and

oxidation at elevated temperature, at a reasonable low cost [16-19]. Ni alloys typically used in

molds protection are divided into two groups: chromium-containing and chromium-free, e.g

NiCrBSi and NiBSi, respectively. The chromium containing alloys are characterized as low

Cr (~8 wt%), medium Cr (~11 wt%) and high Cr (~20 wt%) [20]. The alloys compositions

can be varied in order to optimize specific properties, such as the tribological behavior, the

oxidation resistance, the wettability, etc. Nickel, as the base element of Ni-based alloys, is

responsible for the excellent oxidation and corrosion resistance at high temperature [21, 22].

Other properties can be achieved with alloying elements. B is added to reduce the melting

point of these alloys (concentrations up to 2.5 wt%) favoring the manufacturing ability [22,

23]. Silicon increases the self-fluxing properties, the castability and the wetting characteristics

[22]. Both B and Si act as deoxidizers, forming borosilicate phases, which prevent the

oxidation of the active elements in the alloys [24]. B and C react with chromium to produce

hard borides and carbides and, hence, rising the cavitation and wear resistance of the coatings

[25, 26]. Cr and Al form chromia and alumina which improve the protection of the coatings

against oxidation [27, 28]. The addition of hard and stable carbides and oxides, such as WC,

CrC, SiC, Al2O3, TiO2, ZrO2, CeO2, etc, is used to improve their tribological performance

[29-33]. In conclusion, this wide range of compositions makes these alloys very versatile,

allowing to achieve the desired properties for a particular application.

Ceramic coatings have been extensively applied in components submitted to extremely

high temperature conditions due to their superior properties when compared to metallic

coatings [24-36]. However, their use as thick coatings in the protection of parts for the glass

industry is limited due to their extremely low thermal conductivity. This is a fundamental

parameter for the productivity, which can determine the selection and the applicability of any

solution in glass molds. As it will be referred later, the applicability of ceramic materials can

only be envisaged if this shortcoming is overcome. In order to take advantage of the excellent

high temperature properties of ceramic materials, the only solution is to deposit them in the

form of thin films. An emergent deposition process able to satisfy this requirement is the PVD

technique.

Independently of the base material and the coatings on it applied, the failures in the

molds during service depend on a high number of factors, loading conditions and functional



requirements. Whenever a failure occurs, the most important factor is not the cost of the

component but the costs due to machine down time, process instability and poor production.

Therefore a proper design and fabrication procedure of the molds, as well as their appropriate

use in service, should be carefully analyzed. The factors that lead to the molds failure are the

thermal and mechanical strains, the wear plastic deformation, the thermal fatigue and the

action of the oxidative, abrasive and corrosive environments occurring in service conditions

[37-39]. The synergetic effect of the harsh conditions with a faulty design, a defective

material or bad working practices is the perfect combination to induce the premature failure of

components, expressed as: worn-out geometries, plastic deformation, surface and bulk

material cracks, breakdown of the coatings. For example, the presence of sharp corners,

notches and sudden changes in the cross section should be avoided in molds design. They act

as hot spots which potentiate the failure by thermal fatigue, wear and cracking due to the high

temperature there reached [38, 40, 41]. With the components of molds being constantly

heated and cooled down during glass containers production, large thermal gradients are

created causing the mold tension during heating and compression during cooling down [37,

42]. Thus, surface cracks are frequently produced resulting in a poor product surface

finishing. Moreover, if the thermal expansion coefficients of the base materials and the

coatings are very different, cracks can appear on either base material, base material/coating

interface or, even, only in the coating, leading also to the mold failure. Thermal shock is also

reported as a cause for premature mold failure [37, 38]. In conclusion, the increase of molds

lifetime, the improvement in glass containers quality, the decrease in the rejection rate and the

reduction of costs production can only be achieved through a convenient selection of new

materials (substrate and coatings), the optimization of molds geometries, application of

innovative processes of deposition and adoption of new working practices.

As it was referred to in Chapter I, the main processes used to apply coatings to protect

the mold surfaces in glass industry, are: (i) hardfacing processes such as Plasma Transferred

Arc (PTA) or Gas Tungsten Arc Welding (GTAW), and (ii) thermal spraying processes such

as Flame Spraying (FS) and High Velocity Oxy Fuel (HVOF). Ones more than the others, in

general, these processes often lead to problems of adhesion of the coating to the substrate,

high rate of porosity (5-15%), high residual stresses, excess of dilution with the consequent

changes in microstructural and functional properties of the coatings and base materials, in all

cases compromising the performance of the molds and the final quality of the produced parts

[43-47]. FS is the oldest process used for depositing coatings on molds for the glass industry,

since it is simple, easy to be applied and of low cost when compared with other deposition



methods. However, the adhesion of the coating to the substrate is very low and it presents a

high level of porosity, which frequently reduces significantly the life time of the molds in

service [9]. To avoid adhesion problems, hardfacing processes have been preferentially used

(GTAW and PTA), since they promote a strong metallurgical bond between the coating and

substrate, condition required to avoid their catastrophic failure in service. This technique

gives rise to much better results than thermal spraying processes, for which the adhesion is

only due to a mechanical bonding. GTAW process deposits thicker coatings with lower

porosity and high production rates than flame spraying [48]. However, it is generally

restricted to flat or horizontal surfaces. On the other hand, PTA process produces better

results than GTAW in terms of efficiency and quality of the coatings, at a lower cost. It is a

very versatile process allowing depositing any kind of material, with lower porosity at a

higher production rate than FS and GTAW [49-52]. This process assures a deep penetration of

the coating material into the bulk substrate with a narrow heat affected zone and a relatively

small weld bead due to its high velocity plasma jet when compared to GTAW. However, it is

not portable and cannot be used to coat complex forms, being its main disadvantages.

Presently, PTA is being increasingly used in components for glass industry, mainly on the

mold edges, zone submitted to a high abrasion by the hot plasticized glass. Furthermore,

despite the good adhesion promoted by the hardfacing processes due to the formation of a

metallurgical bond, in some cases this can be a disadvantage. In fact, to reach high adhesion

values it is frequently necessary to increase the level of dilution of the base material with the

coating. Such a procedure leads to changes in the chemical composition of the deposited filler

metal, which can compromise the properties of the coatings and their performance in service.

Hence, a proper selection of the deposition parameters should be carefully done to accomplish

the commitment between the requirements of mechanical strength, corrosion/oxidation

resistance and adhesion [44]. Normally, coatings with thickness in the range of 2 - 5 mm are

applied in the molds surfaces with these processes.

HVOF process provides coatings with higher hardness than PTA do, since it allows

higher cooling rates which promote refined microstructures [9, 53, 54]. Furthermore, owing to

the high energy of the particles jet, better mechanical interlocking can be achieved with the

consequent increase of the bond strength of the coatings to the substrate in relation to FS

process; however, it continues significantly lower than that of the hardfacing processes [55].

As this process is based on the spraying of powders in the melted state over a cooled

substrate, the heat affected zone of the base material remains very low, avoiding the formation

of brittle phases resulting from the metallurgical melting process occurring frequently in



hardfacing processes, which lead very often to failure problems by cracking. The main

disadvantage of this process is the need of fusion of the coating after deposition, since the

rapid cooling down of the powders from the melted state does not allow a good intermixing

among them and a “poor” cohesion of the particles is normally achieved. HVOF allows to

deposit uniform coatings (without dilution) keeping heat extraction conditions uniform in the

entire molds surface, which is a required condition to evenly cool down the glass containers

and, therefore, to produce glass parts without defects. HVOF is being used to coat critical

components in the glass industry, for example the plungers shown on the right in Figure 2.1,

where an efficient and uniform heat transfer is required [2, 3, 8]. The typical thickness of the

coatings for molds surface protection is in the range of 100 - 500 µm. A summary of the main

and typical characteristics of thermal spraying and hardfacing processes is shown in Table

2.1.

Table 2.1 - Comparison of the characteristics of several hardfacing and thermal spraying processes [9, 55-60].

Dep.

process

Material

feed type

Spray gun

temperature

(0C)

Temperature reached in

the electric arc (0C)

Particle

velocity

(m/s)

Relative

bond

strength

Porosity

level

vol.%

Dep. rate

(Kg/h)

FS Powder 3000 - 30-120 Fair 10 - 15 1 - 10

GTAW Wire - 6000-11000 - Excellent 0.01 - 20 6

PTA Powder - 12000-18000 - Excellent 0.1 - 1.5 8

HVOF Powder 4000 - 2000 Good to

Excellent 0.1 - 2 2 - 12

Recently, other deposition techniques, such as Atmospheric Plasma Spraying (APS)

and Physical Vapor Deposition (PVD), have emerged, showing a promising potential for

application in glass industry. APS process is a thermal spraying process very similar to HVOF

since both are based on the projection of melted powders towards the substrate [61, 62].

However, whilst in HVOF process the material in the form of powder is heated based on the

generation of a hot flame obtained by combustion of gases, in APS process melting is

promoted by the powder injection into a very high temperature plasma generated by a

discharge, where it is rapidly heated and accelerated. The main advantage of the APS process

is the much higher temperatures that can be reached in the plasma, making it suitable to

deposited materials of high melting points, as it is the case of cermets and ceramics [59]. In

HVOF the temperature reached is around 4000 ºC and the particle velocity 2000 m/s, while in

plasma spraying the temperature reaches 20000 ºC and the velocity 400 m/s. However, the

very high temperatures and lower velocity occurring in the APS process results in coatings

with microstructures containing oxides, nitrides and, also, some porosity as well as lower



adhesion to the substrate. Despite these disadvantages, APS deposition of cermet and ceramic

coatings was shown in some cases to lead to much better properties than when they are

deposited by HVOF [59]. Thus, the use of the APS process to deposit either the previous

referred cermet coatings (Ni-based alloys reinforced with stable carbides and oxides), which

are normally used to improve the wear behavior, or new cermet coating systems, should be

taken into account. In spite of the extremely high number of studies regarding the effect of

carbide and oxide additions on the tribological behavior of Ni-based alloys, no reports have

been published regarding the effect of nanostructured zirconia additions on their properties.

Presently, pure tetragonal zirconia coatings are one of the main materials used on the turbine

blades surface protection due to their extremely high performance at high temperature [63-

66]. As in glass industry, such material cannot be used as pure coating to protect the mold

surfaces due to its very lower thermal conductivity; however, owing to its high level of

hardness, wear and oxidation resistance, its use as reinforcement material in Ni-based alloy

coatings should be considered.

PVD is a process of deposition of thin films (thickness in the range from few to

thousands nanometers) that produces coatings with outstanding mechanical properties. It is

common in this process, depending on the coating system, to deposit films with two to four

times better mechanical properties than those of the thick coatings described above. The

application of this technique to protect molds for the glass industry is still embryonic, since

their much smaller thickness seems, at industrial eyes, unpromising and unattractive when

compared to that of thick coatings. However, in other industries, where parts are submitted to

extreme high temperature and mechanical loading conditions, these coatings have shown to

successfully increase the life-time and performance of the coated parts, in some cases in more

than one or two orders of magnitude [67-73], allowing predict a potential use of these films to

protect glass molds surfaces. From the development point of view, this technique has been

shown to be very versatile. It easily allows scan the chemical composition of a coatings

system, containing one or more elements, permitting the optimization of the deposition of a

specific film, i.e. the coating with the best compromise between e.g. mechanical, tribological,

oxidation and corrosion properties. Moreover, it allows overcoming the problems of using

ceramic coatings in molds protection. In fact, the lower thermal conductivity of the ceramics

materials can be avoided if the thickness of the non-thermal conductive film is very low. The

main disadvantages of this process are its high investment and production costs, although in

some cases the latter can be decreased. In fact, using PVD process some steps of the molds

production can be removed (for example post-deposition machining tasks, see Table 2.2 for



more detail) and, depending on the mold dimensions, several parts can be coated at the same

time. Therefore, sometimes the final cost of molds production can reach values similar to

those achieved with thick films processes. Finally, it should be taken into account the costs of

glass containers production. On the one hand, the much better properties of thin films in

comparison to thicker ones, may envisage an extended lifetime of the components and

therefore, even if the cost of mold surface protection is higher, the final production costs can

decrease. On the other hand, due to the very low thickness, the possibility of an increased heat

extraction, during glass forming process, gives rise to either important savings in energy,

decrease of the production times or improvement of the final quality of the produced parts.

Table 2.2 - Typical sequential tasks used in molds production.

Deposition processes

FS GTAW PTA HVOF APS PVD

Op

erat

ion

s

Base material machining x x x x x x

Open of U-shaped grooves for

coating deposition x x x - - -

Cleaning x x x x x x

Grit blasting - - - x x -

Coating deposition x x x x x x

Machining x x x x x -

Cleaning x x x x x -

Remelting - - - x yes if cermets -

no if ceramics -

Polishing - - - x x -

In the last decades, the development of thin coatings for mechanical applications was

concentrated in solutions that could ensure simultaneously the specifications required for a

longer lifetime and an increased productivity, high hardness and high toughness, associated

with high thermal resistance, which should accomplish the desired excellent tribological

behavior at high temperatures. The first generation of hard coatings was concentrated in

single carbides, nitrides and oxides (e.g. TiN, WN, TiC, AlN, WC, Al2O3) [67, 74-76]. The

following generations comprised the mixed compounds, such as TiCN, WCN, (TiAl)N,

TiSiN, etc. [77-81] and the multilayer and multiphase structures (TiN/NbN, TiN/AlN, TiN/W,

WN/Ti, TiSiN, TiBN) [80, 82-84]. Finally, super hard coatings were developed as for

example, c-BN, c-B4C, etc [85, 86]. In all these cases, the extension of the lifetime and

performance of the coated components was successfully achieved.

The development process of hard coatings based on transition metal nitrides started

with single nitrides, such as TiN and CrN, which have been widely used as protective



coatings [87-90]. However, the need for using these coatings in more severe and demanding

environments (particularly, at high temperature applications) led to their alloying with

chemical elements that could in anyway improve their thermal behavior; here multi-element

nitride films appeared [91, 92]. The most known of these is TiAlN, which has higher

mechanical strength and oxidation resistance than TiN [93, 94]. The entire range of Ti/Al

ratios was scanned and, globally, the use of very high Al/Ti ratio gives rise to a significant

improvement of the oxidation resistance at high temperatures [95, 96]. With the same aim,

improvement of the oxidation resistance, the addition of Cr to TiN film has also been

extensively studied [92, 97, 98]. Cr was better in some situations, Al or Si were in others.

Therefore, development of quaternary (e.g. TiAlCrN, TiAlSiN, CrAlSiN) and higher (e.g.

TiAlCrSiN) systems, was performed by the incorporation of other elements such as Si, B, Cr,

Al [71, 99, 100]. Among the ternary coatings, TiSiN was one of the most extensively studied

[101, 102]. Depending on the deposition conditions, these coatings have been reported

consisting of: i) nano-sized TiN crystallites surrounded by a Si-N amorphous matrix [103-

105], or ii) substitutional solid solution of Si in TiN structure (not predicted by the Ti-Si-N

phase diagram) [103, 106, 107]. Their main advantage in relation to previous ternary coatings

is the much higher hardness values that can be achieved, between 40 and 50 GPa, for the same

level of oxidation resistance, particularly if a nanocomposite structure is deposited [95, 103,

108].

Currently, during glass containers production, liquid lubricants are used to avoid the

adhesion between the molds surface and the melted glass. In service, most of the lubricant is

volatilized, due to the high temperatures, promoting the glass adhesion to the surface and,

thus, accelerating the wear of the component and giving rise to poor quality in the produced

parts. Furthermore, these lubricants contain commonly small amounts of sulfur which,

combined to oxygen, forms sulfuric acid that erodes the mold surfaces creating holes and

leading to the premature failure of the mold. Thus, the use on mold surfaces of solid

lubricants, instead of the liquid ones, is of great interest, since their low vapor pressure and,

hence, sublimation does not occur so easily. A wide range of solid lubricant coatings have

been developed in the last decades and successfully applied in tribological components. Solid

lubricant coatings such as WC/C, MoS2, diamond-like carbon (DLC), hex-BN, etc, as well as

their combination in nanocrystalline or multilayer structures, are some examples [70, 109,

110]. However, considerable degradation of the tribological effectiveness of these coatings at

high temperature has been reported due to their low oxidation resistance. To overcome this

shortcoming, new concepts of lubrication have been proposed. Solid lubricants for high



temperature applications can be divided into three categories [111-113]: (i) soft metals (e.g.

Ag, Cu, Au, Pb, and In); (ii) fluorides (e.g. CaF2, BaF2, and CeF3); and, (iii) metal oxides (e.g.

V2O5, Ag2Mo2O7). These coatings were developed by combining the intrinsic properties of

some binary or ternary coatings, that are resistant to oxidation, with specific elements in order

to get the lubricious properties.

All the above three types of solid lubricants materials plastically deform and/or form

low-shear-strength surfaces at elevated temperatures, characteristics responsible for the

lubricious effect. Among them, special attention has been given to the coatings based on the

formation of lubricious oxides, being vanadium the preferred element, reason why particular

attention has been given to the vanadium-containing coatings with the expectation of

formation of Magnéli phases VnO3n−1. These coatings showed interesting tribological

properties in the temperature range 500 – 700 ºC [70, 114-119]. Various series of V rich

coatings have been developed, such as ternary (V,Ti)N [120], CrVN multilayered AlN/VN

[121] and quaternary single layer or multilayered CrAlVN [122, 123] and TiAlVN [116, 124-

126]. A detailed description of the effect of V additions on the mechanical, tribological and

oxidation, properties of these coatings can be found in a recent paper review published by

Franz et. al [127]. In summary, the most important conclusions of these studies are: (i) the

relative amounts of V2O5 detected at the oxidized surface of V rich films successfully

decreases the wear rate and friction coefficient of the coatings; (ii) a strong out-diffusion of V

significantly degrades the oxidation resistance of the coatings; (iii) the out-diffusion only

permits a short time efficiency of solid lubrication with the V-oxides.

The control of the V out-diffusion is now one of the major challenges to reach an

adequate oxidation resistance and suitable tribological properties without compromising the

original properties of the host binary and ternary films. One of the possible proposals to solve

this problem will be to use coatings systems composed by a dual phase, in which one of the

phases act as a diffusion barrier to vanadium. This can be the case of TiSiN system. If

deposited as a nanocomposite structure (nano-sized TiN crystallites surrounded by an

amorphous matrix of Si3N4), the Si-N phase may work as an anti-diffusion barrier [128] and,

therefore, if V is incorporated in solid solution in the lattice of the TiN grains, the controlled

release of V can be achieved. By tailoring the nanostructure of the Ti-Si-N films, i.e

producing coatings with different Si-N layer thickness, the V out-diffusion can be controlled

as well as its release as oxide in the coatings surface. An alternative method would be to

deposit a multilayer structure with Si-N layers of increasing thicknesses from the top to the

bottom of the coating.

Chapter III – Effect of the arc current variation on the properties of Ni-based coatings deposited by PTA process


Chapter III

3. Effect of the arc current variation on the properties of

Ni-based coatings deposited by PTA process

3.1. Introduction

This chapter is dedicated to the analysis of the effect of the arc current variation on the

microstructure, hardness, oxidation performance, structural and morphology evolution after

annealing and tribological behavior of Ni-based coatings deposited by PTA process, and is

supported by the publications presented in the Annexes A to E.

The effect of the PTA current variation on the microstructure and hardness of the

coatings is presented in annex A. The microstructural changes in the coatings as well as in the

respective heat affected zones were analyzed in this paper. Further, the effect of a post-weld

heat treatment (PWHT) on the microstructure and hardness of the coatings and base material

was also examined as well as their wear performance evaluated in micro-scale abrasion. This

method was selected to reproduce the main wear mechanism (3 body abrasion) identified in

the surface of molds which have been in service for long time, although it does not allow

consider the high temperature effect. An exhaustive investigation of the test conditions that

produce such kind of wear mechanism was carried out and it is detailed in the paper of Annex

D. The wear resistance of the coatings, in particular their abrasion resistance was evaluated

before and after annealing in order to understand the effect of temperature on the wear



performance of the coatings (publication in Annex C). The influence of the heat treatment on

the hardness and microstructure of the coatings was also discussed. The high temperature

tribological behavior of the coatings was assessed in a pin-on-disc tribometer (publication in

Annex E). Besides the effect of the increasing dilution on the wear behavior of the coatings at

room and increasing temperatures, the base material used in the molds production (gray cast

iron) was also investigated. The publication in Annex B is dedicated to the study of the effect

of substrate dilution on the oxidation behavior of coatings. The oxidation mechanisms

occurring during coatings heat treatment were also discussed in this paper.

The experimental procedure and conditions adopted for the coatings deposition and

characterization are described in detail for each specific investigation on the papers of the

annexes. In the following sections, the main results achieved are summarized.

3.2. Microstructure and hardness of the coatings

Ni-based coatings were deposited by PTA process, using the same process parameters

with exception of the arc current, on flat surfaces of gray cast iron blocks in order to study the

effect of substrate dilution on their properties. The following arc currents were used: 100 A

for specimen 1, 128 A for specimen 2 and 140 A for specimen 4. Further, the edges of a block

with a machined U-shaped groove were coated with the same parameters of specimen 2

(specimen 3).

The effect of the arc current variation on the dilution of the coatings is shown on the

macrographs displayed in Figure 3.1. With the arc current increase in the range 100 - 140 A,

the dilution steadily increased from 28% to 59%. Both flat and U-shaped groove machined

specimens revealed to have similar dilution and macrostructures suggesting that the

deposition conditions on flat surfaces blocks could be extrapolated to molds. As it would be

expected, the increasing dilution had a significant influence on the chemical composition of

the coatings, being the main changes related to the decrease of the Ni content and the increase

of Fe, Si, C and Mo contents, in good agreement with the chemical composition of the base

material.



Figure 3.1 - Dilution induced by PTA process using different weldments: specimen 1 - 100 A, specimen 2 and 3

- 128A, specimen 4 – 140 A.

After deposition, four distinct regions could be observed on the cross section of the

coatings, as follows: the fusion zone (FZ), which corresponds to the coating material (the

nickel alloy), the partially melted zone (PMZ), which is the area close to the FZ where

liquation can occur during weld, the heat affected zone (HAZ), which is not melted but

undergoes microstructural changes, and the base material (BM) which structure remains

unaffected by the deposition process. In all the cases the fusion zone showed a dense

microstructure, free of microcracks and few solidification voids. Increasing substrate dilution

was shown to reduce the content of porosity in the coatings, correlated with the refinement of

the dendritic microstructure (Figure 3.2). The microstructure consisted of dendrites of a Ni-Fe

solid solution phase aligned along the direction of the heat flow. Furthermore, torturous grain

boundaries were displayed, with C flakes (dark-floret structures) evenly distributed in the

matrix. The large number of C flakes on the microstructure was attributed to the dilution of

the base material. Moreover, the dendritic structure became finer as the arc current and,

therefore, dilution increased.

For higher arc currents the heat input during deposition increased, being expected a

coarsened microstructure; however, an opposite behavior was witnessed. According to the

WDS maps (wavelength dispersion spectroscopy) of the most abundant elements of the

coatings, the refinement of the dendritic structure was attributed to the higher number of

precipitates and C-flakes in the grain boundaries, due to substrate dilution. In fact, these

phases normally segregate in the grain boundaries, which increased the rate of heterogeneous

transformation during solidification, hindering the grain growth. A detailed analysis in the

elements distribution with increasing dilution can be confirmed in Annex A.

According to X-ray diffraction patterns of the as-deposited coatings, the major phases

of their structure were primarily (Ni, Fe) face centered cubic solid solution, with the

occurrence of N3Si, Cr5B3 and Fe3Mo3C phases. This phase distribution was in good

agreement with the elemental maps distribution drawn for the coatings. Moreover, according

Specimen 1 Specimen 2 Specimen 3 Specimen 4

28% 50% 54% 59%



to WDS analysis, other precipitates such as C-B, and Mo-C were also observed in the grain

boundaries, contributing to the microstructure refinement, although they were not detected by

XRD analysis, probably due to their small size and/or content. Close to the interface

coating/base material, a cellular microstructure, finer than the dendritic structure, was

observed for all the specimens (see Figure 3.2 d) for specimen 1). This refinement may be

attributed to the high solidification rates involved owing to the efficient thermal exchange

ensured by the high volume ratio substrate/coating.

Figure 3.2 - Optical micrographs of: a) specimen 1, b) specimen 2, c) specimen 4, d) interface of specimen 1.

Despite the previous refinement of the coating microstructure with increasing dilution,

which would suggest an increase of the hardness, an opposite trend was registered, i.e. the

hardness of coatings decreased with increasing arc current, as it is shown in Figure 3.3. This

behavior is undesirable in terms of wear resistance and could only be explained by the

changes in the chemical composition reported above. In all the cases the hardness throughout

the melted material revealed to be approximately constant.

The highest hardness value was observed close to the interface coating/fusion zone,

more specifically on the PMZ (538 HV0.5 for specimen 1), being lower with increasing arc

current, which may be favorable in terms of the toughness of that region. This zone has been

C flakes

Grain boundary

PrecipitatesDeposit material

Base

material

Inte

rfa

ce

Circular island

a)

c) d)

b)



found to be formed by hard cementite and martensite phases which are usually responsible for

the brittle behavior of Fe-C alloys. The hardness of the HAZ was significantly higher than the

base material but lower than the PMZ material, as a consequence of the presence of

martensite and the absence of cementite. The proportion of martensite in the microstructure

decreased with increasing distance from the fusion line and, thus, the hardness dropped down

too.

The microstructure of heat-affected zones and the induced mechanical properties were

directly related to the heat input and, therefore, to the current intensity. The high hardness

values found in the PMZ and HAZ made these regions potentially responsible for many of the

mechanical problems occurring in cast iron welds since they were brittle. Thus, a post-weld

heat treatment (PWHT) at 850 ºC for 1h was performed to the specimens in order to reduce

the presence of brittle phases. As can be observed in Figure 3.4 for specimen 4, the heat

treatment successfully reduced the hardness of the PMZ and HAZ zones down to the base

material values without significant changes of the coating hardness. In this case, PMZ, HAZ

and BM zones were formed basically by a perlitic/ferritic structure, justifying the large scatter

of hardness measurements, as shown in Figure 3.4.

Figure 3.3 - Hardness profile across the interface of PTA specimens.

-4 -2 0 2 4 6 8 10 12

100

150

200

250

300

350

400

450

500

550 Specimen 1

Specimen 2

Specimen 3

Specimen 4

HAZ

PMZ

Base material

Fu

sio

n lin

e

Melted

material

Hard

ness H

V0

.5

Distance across the interface (mm)



Figure 3.4 - Hardness profile across the interface after PWHT in specimen 4.

3.3. Oxidation resistance and diffusion mechanisms

Two of the previous coatings (specimens 1 and 3) were held at 800 ºC and 900 ºC in

air atmosphere during 2 h in a thermogravimetric equipment (TGA), in order to study the

influence of the substrate dilution on the oxidation resistance of the coatings. The surface, as

well as the cross-section, of the oxidized coatings were observed by scanning electron

microscopy (SEM) and analyzed with both energy dispersive X-ray spectroscopy (EDS) and

XRD diffraction. The results are fully presented in Annex B; specimens 1 and 3 were termed

as coatings C100 and C128, respectively. The thermogravimetric analysis revealed that

increasing the substrate dilution the oxidation resistance of coatings decreased, as shown in

Figure 3.5. This behavior was correlated to the chemical composition changes occurred due to

dilution; in particular due to the high iron content introduced in the coating. As a

consequence, different oxide phases were detected in the two coatings, although in each

coating the phases indexed after thermal exposure at 800 and 900 ºC were similar. According

to XRD and SEM-EDS analyses performed at the oxidized surface of the coatings, the oxide

phase sets detected at 800 ºC and 900 ºC were: (i) B2O3, SiO2, Cr5B3 and OAlB for coatings

produced with lower dilution and (ii) Fe2O3 (hematite), Fe3O4 (magnetite), spinels of

NiFe2O4, and small amounts of NiO, B2O3 and B6O for higher dilution coatings. In the latter,

carbon particles were also detected.

-6 -4 -2 0 2 4 6 8 10

100

150

200

250

300

350

400

450

500

550 Specimen 4

Specimen 4 after PWHT

HAZ

PMZ

Base material

Fu

sio

n lin

e

Melted material

Hard

ness H

V0

.5




Figure 3.5 - Isothermal oxidation curves at 800 and 900

0C of the: a) coating deposited using 100 A arc current,

b) coating deposited using 128 A arc current.

Moreover, the analysis performed on the cross section of coatings showed different

types of oxide layers after isothermal oxidation. In the coating with lower dilution the oxide

scale showed to be mainly formed by a protective thick layer of Si-O in which small amounts

of a phase rich in Ni, Si, Al and B (Si-O, Ni-Si and Al-B-O phases) were uniformly

distributed. On the top a B rich oxide was detected too. On the other hand, a dual layer

structure was formed in the case of the coating with high dilution: an external layer of Fe2O3,

with small features of Fe3O4 and NiFe2O4 spinel rich phases, and an internal Fe3O4 layer with

small amounts of an evenly distributed dark phase rich in silicon, nickel and iron. Similar to

specimen 1, small amounts of B-O were also detected on the top. Furthermore, between the

0 1200 2400 3600 4800 6000 72000.0

0.2

0.4

0.6

0.8

1.0

a)

800 ºC

900 ºCW

eig

ht gain

(m

g/c

m2)

Time (s)

Specimen 1

0 1200 2400 3600 4800 6000 72000.0

0.2

0.4

0.6

0.8

1.0

b)

800 ºC

900 ºC

Weig

ht gain

(m

g/c

m2)

Time (s)

Specimen 3



outer and inner layers, a very thin Si-O layer was also identified. Thermogravimetric curves

showed that the coating with lower dilution displayed a parabolic oxidation weight gain as a

function of the time, typical indicating an oxide growth controlled by a protective Si-O layer.

The coating with higher dilution followed the same trend at 800 ºC; however, at 900 0C two

stages in the oxidation curve were detected, the first with a linear increase in mass gain, has

been shown to be a compromise between the outward Fe diffusion through a growing layer of

Fe2O3 and the loss of carbon by decarburization and formation of CO2. The second step

obeyed a parabolic law starting at the moment that the oxide scale, above described, thickened

to a critical value impedding the C liberation, leaving the ion diffusion through the scale the

only mechanism controlling the mass gain. A detailed description of the oxidation resistance

and the diffusion mechanisms occurring with these coatings can be found in Annex B.

3.4. Tribological behavior

The influence of substrate dilution on the tribological performance of coatings was

investigated using two different wear equipments: micro-scale abrasion and pin-on-disc

apparatuses. The first equipment was used since it is among wear tests one of the simplest to

reproduce the wear mechanism closest to what is occurring on the surface of molds in contact

with melted glass (3 body abrasion). However this equipment did not allow high temperature

testing. Therefore, the high temperature tribological behavior of the coatings was assessed in

another tribometer (pin-on-disk), although it did not allow reproducing so well the interaction

conditions between melted glass and the surface of molds.

3.4.1. Micro-scale abrasion tests

In order to use micro-scale abrasion tests to evaluate the effect of substrate dilution on

the abrasive wear resistance of the coatings, an exhaustive investigation of the test conditions

(different abrasive concentrations, loads and sliding distances) that produces 3 body abrasion

wear mechanism was carried out and detailed in Annex D. It was observed that low loads and

higher fraction of abrasive slurry enable 3-body abrasion, whilst, 2-body abrasion becomes

stable at high loads and low volume fraction of abrasive slurry. Moreover, it was perceived

that the specific wear rate essentially depends on the wear rate mechanisms (rolling or

grooving) involved and not on the test conditions employed, since these do not produce

changes in the wear mechanism; i.e. different test conditions which produce similar wear

mechanism can be used to calculate the specific wear rate of coatings. The influence of the



substrate dilution on the micro-abrasion resistance of Ni-based coatings was evaluated in

Annex A and Annex C, using a combination of test conditions that would produce 3-body

abrasion wear mechanism. In both cases the wear results showed that dilution slight decreased

the wear rate of the coatings (see Figure 3.6), in good agreement with the decrease in their

hardness. However, due to the high concentration of abrasive slurry used (103 g per 1×10-4

m3

in Annex C against 34 g per 1×10-4

m3 in Annex A), and owing to the different ball surface

preparation done in Annex C, the wear rate of the specimens revealed to be one order of

magnitude higher. In addition, the wear scars promoted by the tests revealed the presence of

two different wear mechanisms (3-body and 2-body abrasion), although 3-body abrasion was

predominant. Therefore, although a pure 3-body wear mechanism was not produced, as it

would be expected by the results of Annex D, the tests allowed closely reproduced the wear

mechanism present on the molds surface although the high temperature effect had not been

considered.

The heat treatment performed at specimen 4 (at 500 ºC) during 5 and 10 days revealed

not promote changes in their specific wear rate of coatings. However, sample exposed during

20 days exhibited a lower specific wear rate (see Figure 3.6). This was attributed to the

increasing of coating hardness with thermal exposure, due to the precipitation of a Ni-P phase

in the coating matrix. A detailed description on the influence of the annealing treatment on the

properties of coatings can be seen in Annex C.

Figure 3.6 - Specific wear rate of Ni-based coatings in as-deposited and annealed conditions from Annex A and

C, evaluated in a micro-scale abrasion equipment.

0

5

10

15

20

25

30

35

40

45

50

Spec. 4

annealed during

Results from Annex A

Results from Annex C

20 d

ays

10 d

ays

5 da

ys

Spec. 4Spec. 2Spec. 1

Wear

rate

(10

-4 m

m3/N

.m)

Annealing results from Annex C



3.4.2. Pin-on-disc tests

Several weld beads of specimen 1 and 4 (different dilutions – 28 and 59%,

respectively) were deposited in parallel on the flat surface of cast iron blocks, ensuring some

overlapping between them in order to produce coatings large enough to allow perform

tribological tests. The tribological performance of the coatings was evaluated at three different

temperatures: room temperature, 550 and 700 ºC, and compared to the uncoated gray cast

iron. Although in the real working conditions of molds glass acts as an abrasive material, due

to its melted state at high temperature, for the tribological tests we have selected as

counterpart a harder and more resistant material: Al2O3 balls. All the wear tests were

conducted at a constant linear speed of 0.10 m.s-1

, a load of 5 N, a relative humidity of

48±5%, and 5000 cycles. The radius of the wear tracks was set to 5.3 mm.

The tribological tests revealed that at room temperature, the wear rate of the coatings

was independent of the substrate dilution and similar to the gray cast iron (see Figure 3.7). In

all coatings the wear was governed by tribo-oxidation. Similar friction values were measured

in all coatings which were understood by the similar wear debris detected on the worn surface

of the coatings, despite of their discrepancy in hardness and chemical composition due to the

different dilution. Gray cast iron showed a similar level of wear rate as Ni coatings at RT in

spite of its lower hardness (2 times lower than the Ni coatings). This was explained by the

presence of graphite on the worn surface, which acting as a solid lubricant led to lower

friction coefficient and compensated a possible higher wear rate due to the lower hardness.

The wear rate of the coatings abruptly increased when tested at 550 ºC; however, it

dropped with further increase to 700 ºC. Globally, specimen 1 displayed much higher wear

rate than specimen 4 at high temperatures, result that could be interpreted on the basis of the

oxidation resistance of coatings. At 550 ºC, due to the lower oxidation resistance of the

specimen 4, a continuous oxide tribo-layer was formed on their worn surface, protecting it

against wear. In case of specimen 1, due to its higher resistance to oxidation, the oxidation

rate is reduced. Due to the movement of the ball, the destruction of the growing oxide layer is

relatively easier, hindering the formation of a stable oxide layer over the surface and, thus,

leading to high wear volumes due to the direct contact between the alumina ball and the virgin

coating. At 700 ºC, the wear rate of both coatings is reduced due to the formation of

continuous and thick oxide-layers on the worn surfaces. This was not the trend for gray cast

iron for which a monotonous increase of the wear rate with increasing temperature was

observed. In this case, the very low oxidation resistance of the material allowed an easy



removal of the tribo-layer and the depletion of graphite from the wear track. This tribological

behavior is presented and discussed in detail in Annex E.

Figure 3.7 - Variation of the wear rate of the as deposited PTA coatings and the gray cast iron with testing

temperature.

In conclusion, the increase in the arc current, and therefore in the dilution of the

substrate, has detrimental effect on the hardness, oxidation resistance and tribological

behavior of the coatings at room temperature. However, beneficial effect on the high

temperature tribological behavior was observed due to the fast formation of oxide layers

which protect the coating surface against wear. Therefore, the dilution degree should be

optimized in order to get the best compromise between the level of oxidation, wear resistance,

mechanical properties and adhesion to make possible to achieve coatings with the best

performance.

0

50

100

150

200

250

300

350

400

450 RT

550 0C

700 0C

Cast ironSpecimen 4Specimen 1

Wear

rate

(10

-6 m

m3/N

.m)

Chapter IV – Influence of nanostructured ZrO2 additions on the wear resistance of Ni-based alloy coatings deposited by APS


Chapter IV

4. Influence of nanostructured ZrO2 additions on the wear

resistance of Ni-based alloy coatings deposited by APS

4.1. Introduction

This chapter is dedicated to the study of the effect of nanostructured ZrO2 additions on

the microstructure, micro-hardness and wear performance of Ni-based alloy coatings,

deposited by atmospheric plasma spraying (APS) on a low carbon steel, and is supported by

the paper presented in Annex F.

The coatings with nanostructured ZrO2 additions were produced by either powders of

Ni-based alloy and nanostructured zirconia mixed by mechanical alloying or separate powders

using a dual powder injection system available at the APS equipment. In both cases, two

contents of nanostructured zirconia were added to the Ni-based alloy: 20% and 40% in

volume. All the properties of the coatings were compared, interpreted and discussed in

relation to either unalloyed nickel or nanostructured zirconia coatings and based on the

deposition parameters. The coatings produced with powders prepared by mechanical alloying

using 20% and 40 % of nanostructured ZrO2 were denominated as “Ni+20% MA” and

“Ni+40% MA”, respectively, and the coatings deposited using the dual system of powder

injection with 20% and 40 % of nanostructured ZrO2 were designated as “Ni+20% Dual” and

“Ni+40% Dual”, respectively.



4.2. Microstructure and hardness of the coatings

All the coatings showed a typical morphology of plasma sprayed materials, with pores,

lamellae, and partially/un-melted particles (see Figure 4.1). Ni pure coating displayed a

compact and homogeneous microstructure (Figure 4.1 a)), while pure ZrO2 coating showed an

extremely high level of porosity with some cracks, resulting from the tensile residual stresses

generated during cooling down to room temperature (Figure 4.1 b)). Ni+ZrO2 MA coatings

displayed a homogeneous and compact microstructure, with small zirconia particles evenly

distributed in the matrix, whilst Ni+ZrO2 Dual coatings exhibited a porous microstructure, full

of semi-melted Ni powders with large particles of ZrO2 entrapped in their boundaries

suggesting a brittle behavior. Nanostructured zirconia additions progressively increased the

hardness of the coatings. However, coatings deposited using powders prepared by mechanical

alloying displayed much higher hardness than Dual coatings (7.0 - 7.6 against 5.4 - 5.9 GPa,

respectively). This was in good agreement with their lower porosity, higher level of

compactness and finer and more homogeneous distribution of the nanostructured zirconia. In

relation to the structure, the main phases detected corresponded to the ones identified in both

pure Ni-based and ZrO2 coatings, i.e: Ni, Ni-Cr-Fe, Cr23C6, Cr5B3 and ZrO2.



Figure 4.1 - SEM morphology of the: a) Ni-based alloy coating, b) nanostructured ZrO2 coating, c) and d)

Ni+20% Dual and Ni+40% Dual, respectively, e) and f) Ni+20% MA and Ni+40% MA, respectively.


The influence of the nanostructured zirconia addition on the wear behavior of a Ni-

based coating was studied using reciprocating tribological testing equipment with a ball-on-

plate configuration. A harmonic wave generated by an eccentric and rod mechanism imposed

a stroke length of 2.05 mm at a frequency of 1 Hz. A soda-lime glass sphere of 10 mm in

diameter was used as counterpart, as the most suitable material to study the interaction of

glass with the coatings. All the tests were conducted at room temperature during two hours.

Cracksa)

c) d)

b)

e) f)



Four different values of normal load were applied to the coated samples and the volume loss,

specific wear rate and friction coefficient were evaluated.

ZrO2 incorporation successfully decreased the specific wear rate and the friction

coefficient of the Ni-based coating, independently of the deposition procedure adopted (see

Figure 4.2). However, this trend is more accentuated in Ni+ZrO2 MA than in Ni+ZrO2 Dual

coatings. Furthermore, increasing ZrO2 content in MA coatings has a decreasing monotonous

effect on the wear rate whereas an inverse trend was observed in Dual coatings for the highest

ZrO2 content. Several factors, such as the hardness, the microstructure and the wear

mechanisms occurring during sliding, were used to interpret those different trends in the

specific wear rates of coatings, as follows:

(i) Ni+ZrO2 MA coatings showed clean wear tracks with longitudinal scratches

identifying grooving abrasion. Thus, the uniform distribution of small ZrO2 particles,

combined with the higher hardness and toughness of these coatings, impeded the liberation of

large wear debris from the coatings, their plastic deformation and adhesion, leading to low

specific wear rate.

(ii) In the case of Dual coatings, the main wear mechanism observed on the worn

surfaces was of adhesion type, as for pure Ni coating, due to the large areas of nickel exposed

to the counterpart. The semi-melted state of Ni powders, the high level of porosity and the

agglomeration of ZrO2 particles located in-between Ni-based lamellae, made the coatings less

tough, inducing an increase of their wear rate.

(iii) The increase of nanostructured zirconia from 20 to 40% in Dual coatings,

promoted higher levels of brittleness in the Ni lamellae boundaries, increasing the volume

loss of the material.

(iv) Pure ZrO2 coating displayed the highest specific wear rate among all the coatings,

due to its brittle behavior.



Figure 4.2 - Specific wear rate of coatings calculated from the reciprocating tribological tests.

The influence of nanostructure zirconia additions on the abrasion resistance of a Ni-

based coating was also studied. The tests were performed at the most promising coatings (MA

coatings) using the same test conditions as for Ni-based coatings in Annex A, with the

exception of the sliding distances, which were adapted based on the wear mechanism. In fact

for these tests conditions (AISI 52100 steel ball bearing with 25.4 mm in diameter rotating at

a speed of 75 rpm under a load of 0.1 N and an proportion of abrasive slurry of 28% (34 g per

1×10-4

m3 of silica in water)), sliding distances higher than 400 turns showed produce a

mixed-mode wear mechanism (3-body plus 2-body abrasion). Thus, in order to ensure 3-body

abrasion formation the experiments were performed using lower test durations (50, 80, 150,

200 and 300 rpm). The volume loss plotted as a function of the “load × sliding distance”

parameter is plotted in Figure 4.3 and the respective specific wear rate, calculated from the

slope of a straight line fitted to these points, is shown in Figure 4.4. In all cases the main wear

mechanism observed on the wear track was three-body abrasion, in good agreement to the

wear mechanism defined on the surface of molds in real applications. The addition of

nanostructured ZrO2 successfully increased the abrasion resistance of coatings, as above

demonstrated in the reciprocating tribological tests. However, the specific wear rate values

calculated with abrasion tests were two to three orders of magnitude higher than those

achieved in reciprocating sliding tests.

0.0

0.5

1.0

1.5

20

25

30

35

40

Ni +

40%

ZrO

2 M

A

Ni +

20%

ZrO

2 M

A

Ni +

40%

ZrO

2 D

ual

Ni +

20%

ZrO

2 D

ual

Dual ZrO2

Ni

Wear

rate

(10

-5 m

m3/N

.m)

MA



Figure 4.3 - Evolution of the wear behavior as a function of the parameter “normal load × sliding distance”.

Figure 4.4 - Specific wear rate of coatings, calculated from the slope of a straight line fitted to the points plotted

in Figure 4.3.

As a global conclusion, it was demonstrated that nanostructured ZrO2 additions

increased the hardness and improved the tribological behavior of Ni-based coatings.

Furthermore, it was observed that the best properties were achieved when coatings were

produced with mechanically alloyed powders, which promoted an evenly distribution of ZrO2

particles and therefore, an improvement on the mechanical and tribological properties.

0.0 0.5 1.0 1.5 2.0 2.50.0

0.2

0.4

0.6

0.8

1.0

1.2 Ni

Ni + 20% ZrO2 MA

Ni + 40% ZrO2 MA

Volu

me loss (

10

-2 m

m3)

Load x Sliding distance (N.m)

0.0

1.0

2.0

3.0

4.0

Ni + 40%

ZrO2 MA

Ni + 20%

ZrO2 MA

Ni

Wear

rate

(10

-3 m

m3/N

.m)

Chapter V - Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering


Chapter V

5. Study of the influence of V additions on the properties

of TiSi(V)N films deposited by DC reactive magnetron

sputtering

5.1. Introduction

This chapter is on the study of the effect of V incorporation on the chemical

composition, microstructure, mechanical properties, residual stresses, oxidation behavior,

thermal stability and tribological performance of TiSi(V)N thin films, containing different Si

contents, deposited by DC reactive magnetron sputtering. The results presented in this chapter

are supported by the papers in Annexes G to I.

Annex G details the influence of V additions on the chemical composition,

microstructure, mechanical properties (hardness and Young’s Modulus) and oxidation

resistance of TiSi(V)N coatings. The effect of Si incorporation on the properties of the TiN

system was also assessed in this paper, for comparison purposes. The thermal stability and the

diffusion processes occurring during the annealing of TiSiN and TiSi(V)N films are presented

and discussed in Annex H. Annex I is dedicated to the study of the effect of V incorporation

on the tribological behavior of TiSi(V)N coatings at room temperature and its comparison

with that of TiN, TiVN and TiSiN films prepared as references.

Chapter V – Study of the influence of V additions on the properties of TiSi(V)N films deposited by DC reactive magnetron sputtering


5.2. Microstructure, residual stresses and mechanical properties

Three series of TiSiN films with different Si contents, each series with increasing V

content, were deposited in a d.c. reactive magnetron sputtering machine, set with two

rectangular magnetron cathodes working in unbalanced mode. A Ti target, with holes

distributed throughout the preferential erosion zone, and a TiSi2 composite target were used

for the depositions. The different silicon contents were achieved by changing the power

density applied to each target and the V content was varied by changing the number of high

purity rods of Ti and V placed in the holes of the Ti target. Three V contents were added to

the films by using 4, 8 and 12 pellets of V. In all cases the total power applied to the targets

was 1500 W. To serve as reference, a TiN film was deposited from the Ti target using 2000

W of power. Ti and TiN interlayers were deposited as bonding layers, in order to increase the

adhesion of the coatings. The interlayer also contained V when V pellets are placed in the Ti

target. In all the depositions the time was adjusted in order to obtain coatings with

approximately 2.5 µm of total thickness (including interlayers). The total working gas

pressure was kept constant at 0.3 Pa, using approximately 30 and 17 sccm of Ar and N2,

respectively. A summary of the studied coatings and the nomenclature for them adopted (used

in Annex G) is shown in Table 5.1.

Table 5.1 - Coatings deposited and their nomenclature.

Increasing V content

Power (W)

0P V

(0 at.%)

4P V

(~3 at.%)

8P V

(~7 at.%)

12P V

(~11 at.%)

Incr

easi

ng S

i

con

ten

t

Ti 1000 TiSi2 500

S3-0

(TiSiN)

S3-4

(TiSiVN)

S3-8

(TiSiVN)

S3-12

(TiSiVN)

Ti 1200 TiSi2 300

S2-0

(TiSiN)

S2-4

(TiSiVN)

S2-8

(TiSiVN)

S3-12

(TiSiVN)

Ti 1300 TiSi2 200

S1-0

(TiSiN)

S1-4

(TiSiVN)

S1-8

(TiSiVN)

S3-12

(TiSiVN)

Ti 2000 TiN

As it would be expected, the simultaneous decrease in the power density of Ti target

and its increase in TiSi2 target, allowed producing coatings with increasing Si contents. In a

similar way, increasing the number of V pellets allowed to deposit coatings with increasing V

contents. In all cases, coatings with nitrogen contents close to 50 at.% were achieved. All



films showed a columnar micro-structure. As Si content was increased, the columns size and

the surface roughness were firstly reduced and, then, increased for the highest silicon content.

V additions did not change the global type of the cross section morphology.

As a general trend, the structure of the films could be assigned to a fcc NaCl-type

crystalline phase as e.g. TiN (ICDD card 87-0628). The influence of increasing Si content on

the TiN system and the effect of V content on the TiSiN coating with the lowest silicon

content (S1-y series) is shown in Figure 5.1. Concerning TiSiN films, a loss of crystallinity

was observed for the coating with the highest Si content, suggesting an amorphization of the

structure. Furthermore, a shift of the diffraction peaks to higher angles was observed with

increasing Si content, suggesting the formation of a substitutional solid solution. In fact, with

Si in solid solution, its smaller atomic radius promoted the contraction of the TiN lattice and

the consequent shift of XRD peaks for higher angles. Residual stresses could not be the cause

for the peaks shift since they were compressive and similar in all the films (approximately 3

GPa). The formation of a substitutional solid solution was explained by the low deposition

temperature, together with a low substrate ion bombardment, condition that do not provide the

necessary mobility of the arriving species for Si-N phase segregation and consequent

formation of a nanocomposite arrangement, i.e. nanocrystalline TiN grains embedded in a

Si3N4 matrix. V incorporation in TiSiN films also made moving the peaks to higher angles,

behavior once again related to a smaller unit cell due to the replacement of Ti by the smaller

V atoms. In these coatings the compressive residual stresses were also in the range of ~3 to ~4

GPa.

Figure 5.1 - XRD patterns of TiSiN coatings with Si content ranging from 0 to 12.6 at.% and S1-y coatings with

V additions.

35 36 37 38 39 40 41 42 43 44 45

S1-12 (12.0 at.%V)

S1-8 (7.3 at.%V)

S1-4 (1.1 at.%V)

TiN (200)TiN (111)

S2-0 (6.7 at.% Si)

S3-0 (12.6 at.% Si)

S1-0 (3.8 at.% Si)

TiN

Re

lative Inte

nsity

2



The dependence of hardness and Young’s Modulus on the Si and V contents is shown

in Figure 5.2. With Si incorporation, the hardness of the coatings increased up to a maximum

value of 27 GPa (6.7 at.% Si), trend attributed exclusively to a solid solution hardening

mechanism since, as above referred, the level of compressive residual stresses was similar for

all the films. A further increase in the silicon content led to a decrease in the hardness, in good

agreement with the loss of crystallinity of this coating. The Young´s Modulus monotonously

decreased with Si incorporation, which is related to the progressive importance of Si-bonds in

the TiN lattice.

Concerning the effect of V additions, similar evolution was observed for all the Si

compositions (see Figure 5.2 b) for S1-y series). A slight increase in the hardness and

Young´s modulus was observed for lower V contents, followed by a drop for the highest one.

Similar explanation was suggested, the hardness enhancement was attributed to solid solution

hardening since, again, the compressive residual stress values are very similar. The coatings

with the highest Si content displayed the lowest hardness and Young’s modulus among all the

V rich coatings, in good agreement with their low crystallinity.

Figure 5.2 - a) Effect of Si content on the hardness and Young’s modulus for the TiN system. b) Effect of V

content on the hardness and Young’s modulus forcoating S1-0.

0 2 4 6 8 10 12 14

16

20

24

28

32

36

40

44

48

a)

Hard

ness (

GP

a)

Si content (at.%)

200

220

240

260

280

300

320

340

360Y

oung's

Modulu

s (

GP

a)



Figure 5.2 (continued).

In relation to the adhesion of coatings, only two failure modes were observed on the

scratch track of the coatings; the first cracking and the first chipping. Globally all the coatings

showed to be well adherent to the substrate, exhibiting high critical load values. In general V

incorporation did not have significant influence in the critical loads whereas, globally, the

addition of Si gave rise to a small decrease of their values, particularly for Lc1 (see Figure

5.3).

0 2 4 6 8 10 12 14

16

20

24

28

32

36

40

44

48

S1-y coatings

b)

Hard

ness (

GP

a)

V content (at.%)

200

220

240

260

280

300

320

340

360

Young's

Modulu

s (

GP

a)



Figure 5.3 - a) Adhesion critical loads values for the dissimilar coatings. Coatings were scratched as the normal

force was increased linearly from 5 to 70N. Lc2 critical load for TiN coating was > 70 N. b), c) First coating

cracking and first chiping observed on the scratch track of coating S2-4, respectively.

5.3. Oxidation resistance and diffusion mechanisms

The effect of increasing Si and V content on the onset point of oxidation of TiSi(V)N

systems were studied in a TGA equipment. The films were annealed from room temperature

up to 1200 ºC at a ramp temperature of 20 ºC/min. Si incorporation in TiN film strongly

increased the onset of oxidation, being S3-0 sample the most oxidation resistant. On the other

hand, V addition to TiSiN decreased the onset point of oxidation down to a value close to 500

ºC, independently of the Si and V contents, value even lower than for TiN film.

In order to better study the oxidation resistance of the coatings, isothermal annealing at

selected temperatures and times were performed. TiN was annealed at 600 ºC during 30 min,

S2-0 was heat treated at 900 ºC during 1h and finally S2-8 was annealed at three different

temperatures: 550 ºC during 1h, 600 ºC during 30 min and 700 ºC during 10 min. The results

40

45

50

55

60

65

70

75

80

Lc1 C

ritical lo

ad (N

)

00

8

16

24

32

40

48

56

64

72

80

a)

S1

-12

S1

-8

S1

-4

S1

-0

S3-y series

S2-y seriesTiN

S1-y series

Lc2 C

ritical lo

ad (

N)

S2

-0

S2

-8

S2

-4

S2

-12

S3

-0

S3

-8

S3

-4 S3

-12

First chiping (Lc2) First coating cracking (Lc1)

b) c)



are plotted in Figure 5.4. When time and temperature are considered, TiSiN is far the best

resistant to oxidation. Furthermore, it is also confirmed that V incorporation decreased

significantly the oxidation resistance of TiSiN coatings, even to lower values than TiN.

Figure 5.4 - a) Thermogravimetric isothermal analysis of coatings exposed at different temperatures. Cross

section morphology of coating: a) S2-0 oxidized at 900 ºC during 1 h, b) S2-8 oxidized at 600 ºC during 30 min.

The main oxides detected on the surface of the coatings after annealing can be

summarized as follows: i) in TiN and TiSiN films, TiO2 - rutile and anatase, were the main

oxides detected; however, in the last case, SiO2 was also present, localized beneath the TiO2

layer; ii) in TiSi(V)N at 550 and 600 ºC the main oxides were Ti(V)O2 and α-V2O5, the latter

displaying different orientations depending on the isothermal temperature and time; at 700 ºC,

which is a temperature higher than the melting temperature of α-V2O5, β-V2O5 was formed,

resulting from a reduction process of α-V2O5. In situ XRD analysis performed in TiSi(V)N

coatings in the range from 500 – 750 ºC (see Annex H) allowed to demonstrate that the first

0 1200 2400 36000.00

0.05

0.10

0.15

0.20

0.25

S2-8 500 ºC 1 h

S2-0 900 ºC 1 h

TiN 600 ºC 30 min

S2-8 600 ºC 30 min

a)

S2-8 700 ºC 10 minW

eig

ht gain

(m

g/c

m2)

Time (s)

b) c)1 µm

Oxide scale

TiSiN film

Interlayers

Oxide scale

TiSi(V)N film

Interlayers

c)1 µm



oxide being formed was Ti(V)O2 phase. TiSiN film exhibited a typical parabolic oxidation

weight gain as a function of the time. According to STEM elemental maps analysis a double

oxide layer was formed in this case: on top a TiO2 layer followed by a SiO2 layer. Thus, the

excellent oxidation resistance of the TiSiN film was a result of the formation of a continuous

protective SiO2 layer in the interface film/oxide scale (for more detail see Annex H) that

hindered the ions diffusion and, thus, protecting the coating from further oxidation. In the case

of the V rich coating, the isothermal curves tested at 550 and 600 ºC showed two steps: at an

early stage, the weight gain increase over time is linear whilst, in the second stage, a parabolic

evolution can be fitted. For higher temperatures (700 ºC) only the parabolic evolution was

observed but with a strong increase in the mass gain due to the melting of V-O phases.

The structural analysis of the oxide scale when the oxidation temperatures are lower

than the melting point of α-V2O5 showed that a two layer-oxide structure was formed, with a

thick porous inner layer (SiO2 plus Ti(V)O2) and an outer discontinuous layer of α-V2O5. To

the first linear step in the oxidation curves was suggested the growing of Ti(V)O2, as a

conjugation of a fast rate of mass transport in the oxide with a slower surface step

accompanying the gas dissociation, adsorption and subsequent entry of oxygen into the

Ti(V)O2 lattice. In fact, at the very beginning of the oxidation process, due to the high Ti

content, a TiO2 layer started growing. The presence of V ions and its high solubility with Ti

promoted the formation of Ti(V)O2 solid solution, which comprised vanadium cations with

lower oxidation states, as V3+

ions [119, 129]. The substitution of Ti4+

by V3+

ions in the TiO2

lattice would increase the concentration of interstitial metallic Ti4+

ions and, simultaneously,

would decrease the number of excess electrons. The oxidation mechanism was inverted in

relation to Ti-N, being the outwards diffusion of metallic ions (Ti3+

Tin+

) more favorable than

the inwards diffusion of O2-

ions. There were several consequences of this inversion: (i) the

oxidation rate increased substantially, (ii) the kinetics of the oxidation was not controlled any

more by a diffusion process but by the rate of dissociation, adsorption and combination at the

surface of O2-

with metallic ions, and (iii) the continuous and compact Si-rich layer cannot be

formed. α-V2O5 and β-V2O5 phase showed to be to be formed exclusively on the top of the

inner layer due to the V ions reduction on the surface. The parabolic beahaviour after the first

linear increase in mass was explained by the oxide scale thickness increase, making harder the

ion transport. Then, from a threshold thickness on, the ionic diffusion is the controlling step of

oxidation, over the dissociation mechanism, giving rise to the parabolic behavior

characteristic of ion diffusion controlled processes.



A detailed description of the oxidation mechanism occurring in TiSiN and TiSi(V)N

coatings oxidation can be found in Annex H.


The tibological behavior of TiSi(V)N coatings at room temperature was assessed in a

pin on disc equipment using two different balls, alumina (Al2O3) and high speed steel (HSS).

The tests were performed with the same test conditions used for Ni-based coatings deposited

by PTA. Coatings S2-4 and S2-8 were used to study the effect of V additions to the TiSiN

(S2-0) and S3-8 film, which had a higher silicon and similar vanadium contents in relation to

S2-8 coating, was used to understand the role of the Si on the films. Besides TiSiN (S2-0)

film, the properties of the V-containing coatings were compared and discussed in relation to

the TiN and TiVN films. The TiVN coating had a hardness of 23.8 GPa, a Young’s Modulus

of 398 GPa, a similar level of residual stresses and slight higher adhesion critical load than

any of the other coatings.

The wear rate of the coatings tested at room temperature is shown in Figure 5.5. TiSiN

was the coating showing the highest wear rate. With increasing V content the wear rate

progressively dropped. Furthermore, for Al2O3 balls the wear rate was much lower than when

the coatings slided against HSS balls. These results were in good agreement with the lower

friction coefficient achieved when Al2O3 balls were used. The addition of V had the same

influence for the TiN reference coating with a clear improvement of the tribological behavior.

The worst performance of TiSiN among all the tested coatings was attributed to its low

fracture toughness. When TiN and TiSiN coatings were tested against Al2O3 balls, an abrasion

wear mechanism was shown to occur on both wear tracks. However, due to the lower fracture

toughness of the TiSiN, deeper scratches and large loosely wear particles were formed in the

wear track, leading to a strong abrasion of the surface and, thus, a higher wear rate.

The decreased of the wear rate and the friction coefficient with vanadium

incorporation was attributed to the formation of the V2O5 phase on the wear track of the

coatings. This oxide acted as a lubricious tribo-film, protecting the coating from wear, being

the protective effect as efficient as higher is the amount of V2O5 formed. In relation to S2-8

film, TiVN coating (both with similar V content) presented lower wear rate. Such behavior

was explained based on the effect of Si on the inhibition of V2O5 formation, resulting in

higher wear rate and friction coefficient.



When tested against HSS ball, the counterparts were covered with both as-deposited

and oxidized material; however, in TiSiN coating higher amounts of big wear debris were

detected, resulting from the lower fracture toughness of this coating. Therefore, a higher

abrasion of the surface was produced with consequent higher wear rate. For V-containing

coatings, the much higher wear rates and friction values measured with HSS in comparison to

Al2O3 balls were attributed to a less V2O5 oxide formation on the worn surfaces.

A detailed characterization of the effect of V additions on the coatings wear and their

correlation to the main wear mechanisms is shown in Annex I.

Figure 5.5 - Wear rate of the coatings tested against Al2O3 and HSS balls.

The effect of V incorporation on the abrasion resistance of TiSiN coatings (series S2-y

films) was also assessed using the same test conditions as for the Ni-based coatings evaluated

in Annex A. However, in this case, lower sliding distances were selected due to a different

wear mechanism and a much lower coating thickness. The volume loss of coatings plotted as

a function of the load times sliding distance parameter is presented in Figure 5.6; the

respective results for the specific wear rate are plotted in Figure 5.7. As for the other

tribological tests, the addition of vanadium gave rise to lower specific wear rates. This trend

was more notorious for coatings with low V content. In all cases, the main wear mechanism,

suggested by the analysis of the wear scars produced by the ball was 3-body abrasion, which

reproduced the wear observed on the surface of real molds after a period of service. In

comparison to pin-on-disc tests, the specific wear rate achieved with this abrasion test was

higher about two orders of magnitude.

0

2

4

6

8

10

12

14 Wear track

Al2O

3 ball

HSS ball

S3-8S2-8S2-4S2-0TiVNTiN

Wear

rate

(10

-6 m

m3/N

.m)



Figure 5.6 - Evolution of the wear behavior as a function of the parameter “normal load × sliding distance”.

Figure 5.7 - Specific wear rate of coatings, calculated from the slope of a straight line fitted to the points plotted

in Figure 5.6.

In summary, although the nanocomposite structure was not formed in TiSi(V)N films,

V additions significantly improve their mechanical and tribological properties; however, it has

a detrimental effect on their oxidation resistance due to the outwards V ions diffusion to the

surface impeding the formation of a continuous Si-O protective layer. The improvement of

the tribological properties with V incorporations was related to V2O5 formation on the sliding

contact that acted as a lubricious tribo-film, decreasing the friction and protecting the coating

from wear.

0.5 1.0 1.5 2.0 2.50.0

1.0

2.0

3.0

4.0

5.0

6.0

7.0 S2-0

S2-4

S2-8

S2-12

Volu

me loss (

10

-4 m

m3)

Load x Sliding distance (N.m)

0.0

0.5

1.0

1.5

2.0

2.5

3.0

S2-12S2-8S2-4S2-0

Wear

rate

(10

-4 m

m3/N

.m)

Chapter VI - Conclusions and Future Research


Chapter VI

6. Conclusions and Future Research

The main aim of the research activity carried out in this thesis was the improvement of

the surface properties of molds used in glass containers production, through the application of

several types of coatings. Two methodologies were assessed: i) the optimization of

conventional Ni-based coatings currently applied on mold surfaces; ii) the development of

new coating systems deposited by emergent technologies. In the first part, the effect of base

material (cast iron) dilution, induced by the change in arc current, on the properties of Ni-

based coatings deposited by the PTA process was investigated, while in the second part two

different topics were explored. Firstly, the influence of nanostructured ZrO2 additions on the

properties of Ni-based alloy coatings deposited by APS process were studied while, in the

second topic, the influence of V additions on the properties of TiSi(V)N thin films, containing

different Si contents, deposited by DC reactive magnetron sputtering, was evaluated. The

coatings have markedly different morphological features, physical, mechanical, oxidation and

tribology properties depending on the deposition technology and, thus, their direct

comparison becomes sometimes without sense. Nevertheless, although individual

comparisons were already drawn in each chapter, some common aspects among different

coatings deserve comparison, mainly those directly related with the potential industrial

applicability, as it will be described below.



The dilution degree of the substrate promoted by changing the arc current affected all

the functional properties of the Ni-based alloy coatings. Their hardness, oxidation resistance

and tribological behavior were strongly influenced by the micro-structure which consisted

typically in Ni-Fe fcc solid solution phase, with intendendritic phases rich in Cr-B, Ni-Si and

Fe-Mo-C. Although, the number of precipitates in the grain boundaries increased with the

dilution, the hardness of the coatings was reduced due to the growing influence of the softer

material of the substrate. As a consequence, also the abrasion wear resistance, calculated from

tests performed at room temperature with silica slurry, decreased with increasing substrate

dilution. The same trend was observed in another tribological test (pin-on-disc), having been

observed specific wear rates varying as much as two up to three orders of magnitude. At high

temperature, dilution showed a beneficial effect on the wear resistance of coatings, behavior

explained by the decrease of the oxidation resistance due to the incorporation of base material

elements. In fact, due to the lower oxidation resistance of the coatings produced with higher

dilution, an easy and fast formation of a high amount of oxide debris on the worn surface

occurred, which by agglomeration gave rise to a protective tribo-layer improving the wear

resistance of the coating. The post-deposition heat treatment of the coatings induced a

reduction of the hardness restricted to the partially melted and heat affected zones. In

conclusion, the advantages and drawbacks associated with the degree of dilution indicate that

to produce the best and most performing coating compromises should be found between the

mechanical properties, level of oxidation and wear resistance of coatings.

Nanostructured ZrO2 additions to Ni-based alloy coatings had different impact when

mechanically alloyed mixed powders (MA) or separated powders (Dual) were used. MA

powders deposited coatings displayed a homogeneous and compact microstructure with small

zirconia particles evenly distributed in the matrix; contrasting with a porous microstructure,

full of semi-melted Ni powders with large particles of ZrO2 entrapped in grain boundaries,

responsible for a brittle behavior. In both cases, addition of ZrO2 improved the hardness and

wear behavior. The brittleness of the coatings was determinant on the wear resistance;

therefore, the wear rates increased when going from MA powders, through Dual powders up

to pure ZrO2 coatings. All the APS coatings showed higher hardness values than the PTA

deposited Ni-based ones; however, their micro-abrasion resistance was worse, differing one

order of magnitude, fact explained by the low cohesion between powders. In conclusion, for

any possible application of ZrO2-containing coatings in the protection of glass molds, an

additional melting treatment will be required, in order to improve the cohesion and overcome

the brittleness problems.



Si and V additions to TiN and TiSiN systems, respectively, did not change

significantly either the global type of columnar morphology or the diffraction peaks indexed

to a crystalline fcc TiN-type phase. A shift of the diffraction peaks to higher angles was

always observed with either Si or V additions, indicating the placement of those elements in

substitutional solid solution in the TiN lattice. Therefore, the required nanocomposite

structure needed for the control of vanadium diffusion was not deposited. The presence of Si

and V incorporation in solid solution in TiN gave rise to a slight increase in the hardness,

although for the high contents an inverse trend was observed. In comparison to previous thick

coatings, much higher hardness values could be achieved with the thin films, as shown in

Figure 6.1, which is an indication of a possible compensation of their lower thickness for

glass molds protection.

Figure 6.1 - Hardness of coatings studied in the aim of this thesis, deposited by different deposition techniques.

Si and V had an antagonist behavior in relation to oxidation resistance and tribological

behavior. Moreover, contrarily to what could be expected, silicon which improved

significantly the oxidation resistance of TiN system, had a deleterious influence on the wear

resistance. On the other hand, vanadium that promoted an abrupt drop in the oxidation

resistance of the TiSiN coatings presented an excellent tribological behavior. The low

oxidation resistance of TiSi(V)N films was attributed to the rapid V ions diffusion throughout

the oxide scale, which inhibited the formation of a continuous protective silicon oxide layer

which, in case of TiSiN is the responsible for the high oxidation resistance. However, the

02468

1012141618202224262830

PTA Coatings

APS Coatings

PVD Coatings

Hard

ness (

GP

a)

Spe

cim

en 4

Spe

cim

en 1 N

i

Ni+

20%

ZrO2 M

A

Ni+

40%

ZrO2 M

AZrO

2

TiNTiV

N

S2-

0 (T

iSiN

)

S2-

4 (T

iSiV

N)

S2-

8 (T

iSiV

N)

S2-

12 (T

iSiV

N)

2.91.9

4.97 7.6

10.6

23.2 23

.827 2

7.3

27.6

23



vanadium at the oxidized surface was promptly oxidized forming a lubricious and protective

oxide, the V2O5. Therefore, despite the observed decrease in oxidation resistance of films

containing vanadium, this should not be regarded as a disadvantage from the tribological

point of view, as the V2O5 phase detected can act as a solid lubricant. Among all the studied

coatings, V-containing coatings displayed the worst oxidation resistance (see comparison in

Figure 6.2). Although TiSiN coating was the most oxidation resistant coating, even in

comparison with the Ni-based coatings (specimens 1 and 3) well known for applications at

high temperature, the addition of vanadium to TiSiN degraded completely its oxidation

behavior. At temperatures higher than the melting point of the V2O5 phase (~675 ºC), the

integrity of the V rich coatings is loss due to rapid oxidation. This can be a drawback for the

application of TiSi(V)N coatings in molds surface protection. Thus, further studies to control

the V diffusion and, consequently, to improve the oxidation resistance of the coatings, are

required. The objective will be the deposition of these films with different structures, i.e. as a

nanocomposite (TiVN grains embedded in a Si3N4 matrix) or as a multilayer structure (TiVN

layers alternating with Si3N4 layers).

Figure 6.2 - Comparison of the oxidation weight gain of Ni-based coatings deposited by PTA and TiSi(V)N ones

deposited by DC reactive magnetron sputtering.

The tribological experiments performed at room temperature, either by micro-scale

abrasion or pin-on-disc testing revealed that V additions could improve the wear behavior of

TiN and TiSiN films. A comparison of the specific wear rate of all the films calculated from

the tests achieved by both techniques using the same test conditions (see Figures 6.3 and 6.4)

showed that V-containing thin films performed much better than the thicker coatings.

0 1200 2400 36000.00

0.05

0.10

0.15

0.20

0.25 Specimen 1 900 ºC 2 hSpecimen 3 900 ºC 2 h

S2-8 500 ºC 1 h

S2-0 900 ºC 1 h

TiN 600 ºC 30 min

S2-8 600 ºC 30 min

S2-8 700 ºC 10 min

Weig

ht gain

(m

g/c

m2)

Time (s)



Preliminary tribological tests performed by pin-on-disc of a TiSi(V)N film at high

temperature also revealed a better performance in relation to the thicker ones (Ni-based

coatings deposited by PTA). In fact, even for a testing temperature 50 ºC higher (see Figure

6.5), TiSi(V)N film showed close to one or two orders of magnitude lower wear rates than Ni-

based coatings and uncoated gray cast iron, respectively.

In conclusion, although further studies at high temperature are required to better

characterize the tibological behavior of V rich coatings, the preliminary tests showed that

TiSi(V)N films had potential for the protection of mold surfaces, despite their low thickness

in comparison to thick coatings. Nevertheless, it should be remarked that pin-on-disc and

micro-scale abrasion apparatus did not reproduce well the loading conditions to which molds

are submitted. Therefore the development of new tribology equipment is recommended for

reproducing the real working conditions, i.e. tests should simulate abrasion at high

temperature. Only with such equipment it will be possible to have a proper prediction of the

coatings wear during service.

Finally, several remarks should be considered in this comparison between thin and

thick coatings. In a first glance, it could be thought that a higher coating thickness (typical of

PTA and APS coatings) could compensate a lower wear resistance. However, in glass

containers production, many times dimensional tolerances in the molds are just a few

microns, values incompatible with a possible wear out of the total thickness of a thick coating,

increasing that way, the viability of thin films as a solution for molds protection. On the other

hand, the use of thin films also allows overcome the poor thermal conductivity attributed to

the use of thick ceramic coatings. In fact, their lower thickness permits a much more efficient

cooling of the coating and base material in spite of the base thermal conductivity of both types

of coatings could be the same. Therefore, lower power consumption for molds cooling down

can be expected, thereby reducing production costs. Hence, for future work, thermal

conductivity tests are recommended to compare the cooling rate of molds protected by thick

and thin coatings. In relation to maintenance problems, the use of these self-lubricant coatings

can also remedy the drawbacks that normally liquid lubricant brings (corrosion and

disassembling problems), which frequently led to the molds failure and, therefore, to

increasing production costs. In fact solid lubricious oxides are not as reactive as liquid

lubricants are and the global amount used is much lower. Thus, the probability of corrosion is

much lower allowing more uniform, durable and efficient lubrication of the produced parts.

Furthermore, the use of these coatings has a huge advantage from the environmental point of

view, since they are not as harmful for the environment as the liquid ones are. In relation to



the production of the coatings, despite the high production costs of sputtered films, the

reduction of the steps in molds production, such as final machining and polishing (see Table

2.2 in chapter II) and the high number of components that can be coated at the same time,

when compared to the technology used for the production of thicker ones, mitigate the costs

of protection of the molds. With the expectancy of increasing lifetime, due to their superior

mechanical and tribological properties, thin sputtered films should be considered in relation to

thick coatings in the protection of molds for glass industry. Specifically, V rich coatings are

of great relevance and promising for molds protection. Furthermore, its application could be

extended to related areas where high temperature resistance and lubricious properties are

required, such as the protection of machining tools for high-speed cutting and dry machining

processes.

Figure 6.3- Specific wear rate of coatings calculated from the micro-scale abrasion tests.

0.0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

4.0 PTA Coatings

APS Coatings

PVD Coatings

Wear

rate

(10

-3 m

m3/N

.m)

Spe

cim

en 4

Spe

cim

en 1 N

i

Ni+

20%

ZrO2 M

A

Ni+

40%

ZrO2 M

A

Spe

cim

en 2

S2-

0 (T

iSiN

)

S2-

4 (T

iSiV

N)

S2-

8 (T

iSiV

N)

S2-

12 (T

iSiV

N)



Figure 6.4 - Specific wear rate of coatings calculated from the pin-on-disc tests.

Figure 6.5 - Specific wear rate of coatings calculated from the pin-on-disc tests performed at high temperature.

0

1

2

3

4

5

6

7

8

9

10

PTA Coatings

PVD Coatings

Wear

rate

(10

-6 m

m3/N

.m)

Spe

cim

en 4

Spe

cim

en 1

TiNTiV

N

Cas

t iro

n

S2-

0 (T

iSiN

)

S2-

4 (T

iSiV

N)

S2-

8 (T

iSiV

N)

S3-

8 (T

iSiV

N)

0

25

50

75

100

125

150

175

200

225

250

275

Tested at 550 0C

Tested at 600 0C

Wear

rate

(10

-6 m

m3/N

.m)

Spe

cim

en 4

Spe

cim

en 1

Cas

t iro

n

S2-

8 (T

iSiV

N)

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Annex A


Annex A

F. Fernandes, B. Lopes, A. Cavaleiro, A. Ramalho, A. Loureiro, Effect of arc

current on microstructure and wear characteristics of a Ni-based coating deposited

by PTA on gray cast iron, Surface and Coatings Technology, 205 (2011) 4094-

4106.

Surface & Coatings Technology 205 (2011) 4094–4106

Contents lists available at ScienceDirect

Surface & Coatings Technology

j ourna l homepage: www.e lsev ie r.com/ locate /sur fcoat

Effect of arc current on microstructure and wear characteristics of a Ni-based coatingdeposited by PTA on gray cast iron

F. Fernandes a,b,⁎, B. Lopes b, A. Cavaleiro a, A. Ramalho a, A. Loureiro a

a CEMUC, Department of Mechanical Engineering, University of Coimbra, Rua Luís Reis Santos, 3030–788 Coimbra, Portugalb Intermolde, Rua de Leiria, 95, Apartado 103, 2431–902 Marinha Grande, Portugal

⁎ Corresponding author at: CEMUC, Department of Mecof Coimbra, Rua Luís Reis Santos, 3030–788 Coimbra, Porfax: +351 239 790 701.

E-mail address: [email protected] (F. Ferna

0257-8972/$ – see front matter © 2011 Elsevier B.V. Adoi:10.1016/j.surfcoat.2011.03.008

a b s t r a c t
a r t i c l e i n f o
Article history:Received 16 December 2010Accepted in revised form 3 March 2011

Keywords:Plasma transferred arcNi based alloyMicrostructureWear

The plasma transferred arc (PTA) technique is currently used to coat the edges of moulds for the glass industrywith nickel-based hardfacing alloys. However the hardness and wear performance of these coatings aresignificantly affected by the procedure adopted during the deposition of coatings. The aim of the presentinvestigation is to study the effect of arc current on the microstructure, hardness and wear performance of anickel-based hardfacing alloy deposited on gray cast iron, currently used in molds for the glass industry.Microstructure, hardness and wear assessments were used to characterize the coatings. Electron probe microanalysis (EPMA) mapping, scanning electron microscopy/energy dispersive X-ray analysis (SEM/EDAX) andX-ray diffraction (XRD) were used to characterize the microstructure of the deposits. The effect of post weldheat treatment (PWHT) on the microstructure and hardness was also studied. The typical microstructure ofthe coatings consists of dendrites of Ni–Fe, in the FCC solid solution phase, with interdendritic phases rich inCr–B, Ni–Si and Fe–Mo–C. Increasing the arc current reduces the proportion of porosity and hardness of thecoatings andmodifies their composition due to the increasing dilution of the cast iron. The partial melted zone(PMZ) had a typical white cast iron plus martensite microstructure, while the heat affected zone (HAZ) hadonly a martensite structure. The wear tests showed decreasing wear resistance with decreasing hardness ofthe coatings. PWHT reduces the hardness of the PMZ and HAZ but does not significantly alter the hardness ofthe bulk coating.

hanical Engineering, Universitytugal. Tel.: +351 239 790 700;

ndes).

ll rights reserved.

© 2011 Elsevier B.V. All rights reserved.

1. Introduction

Traditionally, molds for the glass industry are manufactured fromgray cast iron or copper alloys due to their superior thermal behavior,relatively low cost and excellent thermal conductivity. This isassociated with high hardness and good wear resistance and leadsto a high production rate of glass parts. However, the edges of themolds made of these materials are sensitive to wear in heavy dutycycles; so they need to be coated with materials that are moreresistant to high temperature and wear [1].

The coating of these edges is frequently done by plasma transferredarc (PTA), using nickel-based fillermetals [2]. Normally the PTA processproduces very high quality, thicker deposits, offering optimal protectionwithminimal thermal distortion of the parts, low environmental impactand high deposition rates in single layer deposits [3]. It allows perfectlycontrolled deposition of alloys on mechanical parts that are subject toharsh environments, significantly extending their service life [4]. Thecomposition and properties of the coating are greatly influenced by the

type and dilution of the substrate. Balasubramanian et al. [5] suggestthat the dilution of coatings should be controlled through the relevantprocess parameters (transferred arc current, travel speed, powder feedrate, torchoscillation frequency and stand-off distance). Takano et al. [6]studied the influence of arc current variation, arc constriction andplasma gas flow on the characteristics of Co coatings. They concludedthat current intensity is the parameter thatmost significantly affects thecharacteristics of the deposits as higher currents increase the dilution ofthe base material and decrease the hardness of the coatings. Díaz et al.[7] mentioned that the dilution also increases with gas plasma flow instellite 6 coatings.

Nickel-based hardfacing alloys have become increasingly popularin recent years owing to their excellent performance in environmentswhere abrasion, corrosion and elevated temperature are factors. Themicrostructure of Ni-based hardfacing alloy deposits has been studiedusing various alloy compositions and different substrates [8–10].However, most of the studies consider the deposition of nickel alloyson stainless steels; there are no references to characterization of suchdeposits applied to gray cast iron.

Gurumoorthy et al. [8], who worked with nickel-base superalloysdeposited by PTA on stainless steel, reported that the microstructure ofthe deposit consists of γ-Ni solid solution dendrites, carbides, andinterdendritic eutectics composed ofγ-Ni and other phases identified as

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Table 2Deposition parameters for PTA weld surfacing.

Specimen1

Specimens 2 and 3* Specimen4

Main arc current (A) 100 128 140Powder feed rate (rpm) 20 20 20Travel speed (mm/s) 2 2 2Powder feed gas flow rate (l/min) 2 2 2Plasma gas flow rate (l/min) 2.2 2.2 2.2Shielding gas flow rate (l/min) 20 20 20Torch work distance (mm) 13 13 13Oscillation (mm) 4 4 4Preheat temperature °C 480–500 480–500 480–500

*See text.

4095F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106

being Ni-rich borides. The microstructure and mechanical properties ofNi base alloys are greatly influenced by the alloying elements. Forexample aluminum, titanium and niobium are added to strengthen thematerial through the formation of γ′ gamma prime (Ni3(Al,Ti)) andboron and zirconium are added to improve the creep strength andductility [9,10].

This paper details an investigation of themicrostructure, hardness and3-body abrasion behavior of a nickel-based hardfacing alloy deposited byPTA on gray cast iron with different weld currents. Microstructuralchanges in theHAZaswell as the effect of thepostweldheat treatment onthe microstructure and hardness of the coatings are also examined.

2. Experimental procedure

ThecoatingsweredepositedbyPTAtechnology (usingaCommersaldGroup ROBO 90machine) on flat surfaces of gray cast iron blocks using anickel-based hardfacing alloy. This technique obtains thicker coatingswith lower porosity and higher production rate than the Flame Sprayingand Gas Tungsten Arc Welding processes. The composition of base anddeposit materials is given in Table 1. The same process parameters wereused for eachdepositionwith theexceptionof the arc current,whichwas100 A for specimen 1, 128 A for specimen 2 and 140 A for specimen 4.The edges of a mold made of the samematerial but with a machined U-shaped groove were coated with the same alloy using the parameters ofspecimen 2. This specimen is referred to as specimen 3. The surfacingparameters used in the tests are listed in Table 2. The blocks of basematerial were induction heated before coating, in order to reducethermal shock, thus reducing the tendency to crack. After coating, theblocks were left to cool to room temperature in still air. Samplescontaining coating andbasematerialwere removed fromeachblock andthe mold for further analysis.

The samples for microstructural analysis were polished and etchedfollowing conventional procedures. Etchingwas performedusingequalparts of nitric acid and glacial acetic acid, to reveal themicrostructure ofthe deposits. Marble's solution (HCl (50 ml)+CuSO4 (10 g)+H20(50 ml)) was used to reveal the microstructure of the cast iron. Themicrostructure of the different regions of the samples was characterizedby optical microscopy (OP), scanning electron microscopy (SEM), X-raydiffraction (XRD) and energy dispersive X-ray analysis (EDAX). Thechemical composition was evaluated by Electron Probe Microanalysis(EPMA). The hardness profile across the interface of the coatings wasdetermined with Vickers testing using a 5 N load at 15 s.

Some specimens were subjected to a subsequent heat-treatmentat 850 °C for 1 h and then cooled in the furnace to room temperature,in order to reduce brittle phases present in the partially melted andheat affected zones. Microstructural and hardness profiles of thesespecimens were obtained.

Wear tests on the coated surfaces were carried out at roomtemperature in ball-cratering devices. In these devices, a ball is placedon a flat specimen in the presence of abrasive slurry and rotated whilebeing submitted to a specific load in order to produce a spherical cup-shaped depression. This process was used in order to reproduce thewear mechanism present at the surface of the coatings on glass moldswhich have been in service for a long time. An AISI 52100 steel ballbearing 25.4 mm in diameter was used at constant rotational speed of75 rpm. High purity SS40 silica with angular particles averaging

Table 1Nominal chemical composition (wt.%) of substrate and hardfacing alloy.

Base material C Mn Si P S

Grey cast iron 3.60 0.60 2.00 b0.20 b0.04

Hardfacing alloy C Cr Si

Ni-alloy 0.14 2.45 2.56

3.10 μm in size was used as an abrasive to replicate the effect of themelted glass. The slurry used was prepared with a proportion of28 vol.% of silica in water.

The duration of the test, measured by the number of rotations, wasselected according to the material and the test conditions employed.The normal load used was 0.1 N.

Photographs of the spherical depressions were taken in order tomeasure their dimensions. The wear volumewas calculated by Eq (1),where R is the ball radius and b is the crater chordal diameter.

V = π × b4� �

= 64 × Rð Þ ð1Þ

Archard's model, given by Eq (2), is the onemost commonly appliedto wear results to determine a parameter to quantify the wear behaviorofmaterials. In this equation P is the normal load, l is the sliding distanceand k is the specific wear rate. The specific wear rate is the parameterwhich quantifies the wear behavior.

V = k × P × l ð2Þ

To quantify the specific wear rate a linear version of Archard'smodel was applied and statistical analysis was used to estimate theerror of measurement [11].

3. Results and discussion

3.1. Microstructure

Four distinct regionswere observed in the cross section of each of thecoatingsdepositedwith thedifferentwelding conditions, as illustrated inFig. 1 A) for sample 3. Those regions are the fusion zone (FZ), composedof the Ni-alloy and the base material melted during the process, thepartially melted zone (PMZ), which is the area immediately outside theFZ where liquation can occur during welding, the heat-affected zone(HAZ), which was not melted but underwent microstructural changesand the base material (BM), whose microstructure remains unaffectedduring the coating deposition thermal cycle. The same regions wereobserved by M. Pouranvari [12], who studied the welding of gray castiron using nickel filler metal deposited by shielded metal arc welding.

Cr Ni Mo V Ti Fe

b0.20 b0.50 0.50 0.10 0.20 Balance

B Fe Al F Co Ni

0.86 1.08 1.30 0.01 0.08 Balance

Fig. 1. (A) microstructure of various regions in gray cast iron weld in as-weld condition, (B) partially melted zone, (C) heat affected zone, (D) transition heat affected zone–basematerial, (E) base material.

Fig. 2. Dilution induced by PTA process using different weldments: specimen 1 — 100 A, specimens 2 and 3 — 128A, specimen 4 — 140 A.

4096 F. Fernandes et al. / Surface & Coatings Technology 205 (2011) 4094–4106

image of Fig.�2

Table 3EPMA punctual examination of the deposit and base material of specimen 3.

Elements (% W)

Al Si Ti Cr Fe Ni Mo

1 1.00 2.52 0.00 2.36 27.37 56.12 0.23 Coating2 0.96 3.06 0.00 2.50 26.48 56.11 0.173 0.14 0.79 0.06 4.64 16.46 62.24 0.244 0.01 2.17 0.00 0.03 91.73 0.07 0.27 Interface5 0.02 2.05 0.01 0.03 93.81 0.10 0.32 Base material


Fig. 1 B) to E) illustrates the microstructures of the zones marked withthe digits 1 to 4 in Fig. 1 A) in higher magnification.

Fig. 1 B) shows the transition zone between the coating and thepartially melted zone. Fig. 1 C) illustrates that the HAZ is basicallycomposed of a martensitic structure (needle shape), which is foundmainly in the region closer to the fusion line, however, precipitates andgraphite flakes are also present. Themicrostructure of the basematerial isprincipally ferritic with flakes of graphite and precipitates uniformlydistributed in the matrix (Fig. 1 E)). Fig. 1 D) illustrates the transitionbetween the HAZ and the base material. The nature and relative size ofthose zones are determined by the coating procedure, mainly the heatinput during the process, and the composition of the cast iron andhardfacing alloyused. The constitutionof themain zoneswill be discussedseparately, in order to provide a better understanding of the effect ofcurrent on their characteristics.

3.1.1. Melted zoneFig. 2 illustrates the effect of the current on the dilution produced

in each sample. The dilution is defined as the proportion between thearea of melted base material and the total area of the melted zone,

Fig. 3. Aspect of the coatings of: A) specimen 1, B) specimen

both measured from the cross section of the coating using Photoshopimage analysis. These images show that as the current is increased inthe range 100–140 A, the dilution steadily increases from 28% to 59%.Specimen 3 in the image illustrates the groove shape usually machinedin the mold before coating. Comparing specimens 2 and 3, which wereboth coated using the same current, it is possible to see that the grooveproduces only a small change in the dilution and the results obtainedfrom the flat blocks can be extrapolated to molds.

The dilution of cast iron modifies the composition of the coatings,as illustrated in Table 3, which shows the chemical compositionmeasured by EPMA at various points of the coating and the substrateof specimen 3. Point 1 is located close to the coating surface, point 2 atthe mid-thickness of the coating, point 3 in the coating close to theinterface and the other two points are in the bulk cast iron; point 4close to the interface and point 5 distant from it. The amount of iron inthe coating (N16%) is much higher than that provided by thehardfacing alloy (1.08%), producing a consequent reduction in thenickel content. According to Ezugwu et al. [9] the increase in ironcontent in the coatings tends to decrease their oxidation resistancebecause of the formation of a less adherent oxide scale. The increase inMo content in the coating is also caused by the dilution of the cast ironbecause it is absent in the hardfacing alloy. Although the true carboncontent value could be skewed by contamination, carbon is present inthe coating in large quantities, again due to the contribution of thegraphite from the cast iron. Thus, increasing the arc current willincrease the dilution and the chemical composition of the coating willchange accordingly.

All coatings contained small random pores throughout the crosssection, as shown in Fig. 3. According to F. A. L. Dullien [13], the“volumetric” and “area” porosity proportions are equivalent for random-structured porous media, which agrees with our results. The area of the

2, C) specimen 4, and D) fusion interface of specimen 1.

image of Fig.�3

Fig. 4. Optical micrograph of: A) specimen 1, B) specimen 2, C) specimen 4, and D) interface specimen 1.


poresvisibleunder theopticalmicroscopewascalculatedusing the ImageJprogram (developed at the National Institutes of Health, United States).The proportion of porosity calculated is very low for all specimens; 0.35%for specimen 1 with an average pore size of 1.45±0.58 μm, 0.32% forspecimen 2 with an average pore size of 1.61±0.65 μm and 0.28% forspecimen 4 with an average pore size of 1.37±0.40 μm. This decrease inthe proportion of micropores in the coatings as current increases can becorrelatedwith the refinement of thedendriticmicrostructurementionedbelow. The decrease in dendrite size reduces the amount of liquid metaltrapped by dendrites during solidification, reducing the number of voidsin the microstructure. Charmeux el al [14] mentioned this mechanism inthe solidification of Almicro castings.On the other hand the cast iron alsoexhibits significant porosity, with pores reaching 17 μm in diameter, asillustrated in Fig. 3 D).

The typical microstructure of the coatings is shown in Fig. 4. Themicrostructures consist of dendrites of the Ni–Fe solid solution phase,with columnar morphology oriented along the direction of heat flowand torturous grain boundaries. The microstructure contains C flakes(dark floret-like structures) equally distributed through the micro-structure. The chemical composition of the Ni-based powder (lowcarbon content see Table 1) does not explain the large number ofthese flakes which can only be attributed to the dilution of cast ironinduced by the PTA process. Furthermore, small interdendriticprecipitates can also be detected. These are possibly carbides, asexplained below. Fig. 4 A–C) shows that an increase in the arc currentand, therefore, an increase in dilution, allows carbides and C-flakes inthe microstructure to be more easily detected and observed. It wasthought that there could be some coarsening of the microstructure as

the current was increased due to greater heat input. However this wasnot observed. Fig. 4 shows that the dendritic structure becomes fineras the current increases. For example, the average ternary dendritespacing for the specimens 1 and 4 is 15.4 and 11.2 μm, respectively. Asexplained below, this refinement can be related to the change in thecomposition of the deposited material due to its increasing dilution asthe arc current increases.

The phenomenon of circular “islands”, which frequently occur close tothe interface of the coating, is common to all the specimens (see Fig. 4 D)for specimen 1), These islands are characterized by the total absence ofgraphite. According to the Commersald company [15] they are causedwhen powder grains not melted by the plasma arc subsequently melt incontact with the pool. Near the interface a cellular microstructure finerthan thedendritic structurewasobserved for all the samples, as illustratedin Fig. 4 D) for specimen 1, where the grain size is approximately 12 μm.This refinement may be due to the higher solidification rates involved inthe boundary because of the efficient thermal exchange ensured by thehigh volume ratio substrate/coating of the base material, as suggested byGatto et al. [3].

SEM analysis gives amore detailed view of themicrostructure of thecoatings. Fig. 5 shows a SEM image of the microstructure of the coatingof specimen 1. The image shows that in addition to the dendrites of theNi Fe solid solution phase and graphite flakes the grain boundary isformed of two phases: a light gray phase (like-grain) and a dark grayone. EDAX spectra showed that the light gray phase and the dark grayphase are rich in silicon and chromium, respectively.

A detailed analysis was carried out to study themain constituents ofthe coatings using EPMA. WDS (wavelength dispersion spectroscopy)

image of Fig.�4

Fig. 5. A) Typical SEM micrograph (etched) revealing the grain boundaries of specimen 1; SEM EDAX spectra of: B) dark gray phase, and C) light gray phase.


mapsof themost abundant elements of the coating for all the specimenswere obtained. Fig. 6 A) represents a SEM image of themicrostructure ofthe deposit for the specimen 1. The respective elemental maps of Fe, Ni,Si, Cr, Al, B, C andMo are shown in Fig. 6 B–I), color-coded such that thelowest concentration of the element analyzed is indicated in purple andthe highest in red. The results of the WDS maps for specimen 1 revealthat the dendrites are rich in iron, silicon and aluminum, whileboundaries are rich in boron, silicon and chromium.Nickel is distributedalmost uniformly throughout the structure, except in areas wheregraphite flakes are present. In the boundary chromium appears, toexclude siliconas these elements do not overlap. However, chromiumand boron do overlap in the boundaries, as illustrated in Fig. 6 E) andH).Thus, it can be concluded that the borders of the boundary are rich in Siwhereas themiddlesof theboundaries aremainly composedof B andCr.Comparing the information from Figs. 5 and 6 it is possible to conclude

that the light gray (like-grain) phase is rich in silicon and the dark grayphase is rich in chromium,which agreeswith the EDAX analysis. Carbonconcentrates chiefly in the graphite flakes, as shown in Fig. 6 F), thoughat someparticular points the increment in theC signalwasalsoobservedin the boundary, mainly coincidingwith the B signal (see circled zone inB andCmaps) suggesting that carbidesmayhaveprecipitated. Althoughnot visible in the C-map, very small agglomerates ofMo can be detectedin Fig. 6 also suggesting formationof carbides.However, as shownbelowby XRD analysis, these carbides were not identified, probably due totheir small size and/or content.

The increase in the arc current and, consequently, in thedilutionof castiron brought some changes in the distribution of themain elements in themicrostructure, as shown by theWDSmaps for specimen 4, illustrated inFig. 7. For this specimen, in comparison to sample 1, themain difference isrelated to the Ni and Fe signals. It is clear from Fig. 7 that the zones in the

image of Fig.�5

Fig. 6. WDS maps of the coating of specimen 1: (A) SEM image of the microstructure of the deposit, (B) iron, (C) nickel, (D) silicon, (E) chromium, (F) carbon, (G) aluminum, (H)boron, and (I) molybdenum.


interdendritic boundaries richer in iron are depleted in nickel and siliconbut that in the boundary these zones overlapwith Cr and B. The groupingof thesemaps suggests the presence of a Fe–Cr–Bphase in the boundaries,however, as can be seen later from the XRD analysis this is not identified,as only a preferential associationwith the B–Cr phase is indicated, leadingfree iron to combine with other elements. The Si rich zones, see Fig. 7 D),are depleted of Fe and rich inNi. In specimen4, aswell as the strong signalgiven by the C flakes tiny molybdenum and carbon traces could bedetected in the surroundingmaterial (whereC is hardlypresent),which isa clear contrast with the results from specimen 1. These elements aredistributed preferentially in the boundaries overlapping the Fe map,suggesting the formation of a mixed carbide containing Fe and Mo. Theincrease in the content of these elements is caused by the increased

dilution, since they are present only in the cast iron. Collins and Lippold[16] and Ramirez et al. [17] mention that the presence of precipitates inthe interdendritic regions results in the formation of very tortuous grainboundaries and theexcess of theses precipitates stops themigrationof thegrain boundaries. This can justify the lower grain size of the coatingdepositedwith 140 A in spite of the higher energy supplied to the coatingprocess. The literature also indicates that an increase in the molybdenumand carbon content improves the rate of heterogeneous transformationduring solidification. These elements normally segregate in the grainboundaries, thereby hindering grain growth [18]. Nevertheless, theinfluence of the much higher number and content of graphite flakespresented in the 140 A sample, these being segregated from the grainboundaries during material solidification, should therefore also be

image of Fig.�6

Fig. 7. WDS maps of the coating of specimen 4: (A) SEM image of the microstructure of the deposit, (B) iron, (C) nickel, (D) silicon, (E) chromium, (F) carbon, (G) boron, and (H)molybdenum.


considered as anexplanation for the lowergrain size in this sample. As canbe seen in Fig. 4, C flakes are common in the grain boundaries.

The distribution of the chemical elements in the microstructure inspecimens 2 and 3 follows the trend mentioned above.

Fig. 8 shows the XRD spectra of specimens 1, 2 and 4. The analysisof this data revealed that the major phase present in the coatings is a(Ni, Fe) solid solution face centered cubic (ICDD card 47–1405-(111),(200), (220), (311), and (222) peaks at 2θ~51.1°, 59.8°, 89.6°, 111.42°and 119.3°, respectively). WDSmaps show that in Cr and B rich zones,Si does not exist as it is connected to Ni. The overlapping of Ni and Sisignals in WDS maps suggests the presence of a Ni–Si phase, such asNi3Si (ICDD card 32–0699) shown in Fig. 8, corresponding to some of

the lower intensity experimental XRD peaks. The other peaks can beadjusted either to a Cr-boride phase Cr5B3 (ICDD card 32–0278) or tothe Fe, Mo mixed carbide of M6C (Fe3Mo3C — ICDD 47–1191) type.The increasing content of this latter phase through samples 1 to 4 sitswell with the WDS maps presented and interpreted above. Further-more, the indexation is in agreement with the results from theliterature where similar phases were also detected in a nickel alloydeposit on an austenitic stainless steel [8]. In summary, the increase inthe arc current increased the dilution of the cast iron, raising the iron,carbon and molybdenum content in the coating, promoting dendriterefinement of the microstructure due to the increase in precipitates(Fe–Mo–C) and C flakes in the grain boundaries.

image of Fig.�7

Fig. 8. X-ray diffraction pattern of the deposit material obtained for the specimens 1, 2 and 4.


3.1.2. Partially melted zone (PMZ) and heat affected zone (HAZ)These zones are critical in cast irons since the material can solidify

as white iron if cooling is rapid enough. If the amount of graphitedissolved during welding is high enough it is likely that it will alsogive rise to a continuous carbide network. This is undesirable as abrittle carbide matrix can cause durability problems in the final part.Fig. 9 shows the microstructure of this zone in specimen 3. In thisimage, cementite (Fe3C—white phase) can be seen concentrated nearthe interface. Fig. 9 B) shows the microstructure of the PMZ zoneunder high magnification. The image displays a large amount ofcementite in the grain boundaries with acicular martensite inside thegrains. Some precipitates, identified as mixed titanium and molybde-nium carbides (marked respectively with the numbers 1 and 2 inFig. 9 and identified in Fig. 10, were also found. As these phases arehard and brittle this region is sensitive to in-service crack initiationand propagation due to thermal and mechanical fatigue, as discussedbelow. M. Pouranvari [12] also reported that a continuous brittlenetwork of coarse carbides along the weld fusion line may lead toinitiation of cracking.

In the heat affected zone, carbon can diffuse into the austeniteduring welding and the austenite may subsequently transform intobrittle martensite due to the high cooling rate. Martensite is also

Fig. 9. SEM image of partial melted zone and heat affected zone of specimen 3, (A

susceptible to cracking. The amount of martensite formed depends onthe composition of the cast iron and the thermal history of the zone.

3.2. Hardness

Fig. 11 shows the hardness profiles across the interface for all thestudied specimens. The vertical line in the graph represents the fusionline. The hardness through the melted material is approximatelyconstant and tends to decrease with increasing weld current. This canbe related to the higher dilution induced by the process.

For all the samples, the highest hardness valuesweremeasured in thearea near the fusion line. A maximum hardness of 538 HV was observedin specimen 1 and decreases with increasing arc current for the otherspecimens. As previously documented, themicrostructure in that regionconsistedofhardmartensite andcementite. Theproportionofmartensitein themicrostructure decreases with increasing distance from the fusionline and, thus, the hardness decreases too. Themicrostructure in the HAZis directly related to the heat input in the process and therefore to thecurrent intensity used.

The high hardness values found in the PMZ and HAZ make theseregions potentially responsible for many of the mechanical problemsexperienced in welds in cast iron. The most effective way to reduce

) transition between coating and the substrate, and (B) magnification of PMZ.

image of Fig.�8

image of Fig.�9

Fig. 10. Energy dispersive X-ray analysis (EDAX) of the particles shown in Fig. 9 b), (1) titanium carbide with square shape, and (2) molybdenum carbide with rounded shape.

Fig. 11. Hardness profile across the interface.

Fig. 12. Hardness profile across the interface after PWHT in specimen 4.


image of Fig.�10

image of Fig.�11

image of Fig.�12


hardness and cracking problems in these regions is to reduce theseverity of the thermal cycle produced by the process. This can be doneby controlling the heat input and preheating during coating. Anotherway to reduce the hardness in these zones is to submit the parts to anannealing treatment, although this solution is expensive.

The precipitates, the porosity and the random distribution of graphiteflakes in the cast ironmatrix can be linked to the variability of hardness inthe base material.

Fig. 14. Evolution of the wear behavior as a function of the parameter “normalload×sliding distance”.

3.3. Effect of post-weld heat treatment (PWHT) on the microstructureand hardness

Although the heat treatment did not induce significant changes inthe microstructure and hardness of the coatings, it induced substan-tial changes in the PMZ and HAZ, as illustrated by the hardnessprofiles before and after PWHT shown in Fig. 12 for specimen 4. Thisspecimen was produced with the highest arc current, and had thelowest hardness values in the PMZ, as shown in Fig. 11. Fig. 12 showsthat even in this case the heat treatment (holding at 850 °C for 1 h)produced little change in the coating but a substantial decrease in thehardness of the PMZ and HAZ of the specimen. Similar behavior wasobserved for the other specimens. This reduction in hardness can beattributed to the reduction in the proportion of hard phases,cementite and martensite, in these zones. After PWHT the basematerial displays basically a perlitic/ferritic structure.

Fig. 13. Indentations in base material (specimen 4) after PWHT. A) indentation made inthe perlitic structure, and B) indentation made in the ferritic structure.

After heat treatment a large scatter of hardness measurementscontinues to be observed in the cast iron. This scatter can be attributedto local alterations in the microstructure, as shown in Fig. 13, whichrepresents the aspect of the indentations marked with the numbers 1and 2 in the hardness profile in Fig. 12. The image reveals that if theindentation impacts mainly the perlitic structure, the white zone inFig. 13 A), the hardness will be greater than if the indentation isperformed in the ferritic structure, the gray zone in Fig. 13 B). Inconclusion, PWHT is beneficial because it reduces hardness in the PMZandHAZ,without producing significant alterations in themelted zone.

3.4. Wear test

Thewear tests were performed to predict the effect of the differentprocess parameters on the in-service wear behavior of the coatedcomponents. The ball-cratering test results are shown in Fig. 14, as thevolume of material loss plotted against the product of sliding distance(l) and normal load (P). A complete reliability analysis is summarizedin Table 4. The wear volume displays linear evolution with P×l, andthe slope corresponds to the specific wear rate. Specimen 1 exhibitsboth the lowest wear volumes and the lowest specific wear rate,therefore, it is the most resistant to wear. This is compatible with thegreatest hardness displayed by the coating of this specimen, asillustrated in Fig. 11. Samples 2 and 4 display very similar wearvolumes, which is confirmed by similar hardness values; even so,sample 2 has a specific wear rate slightly higher than sample 4.

Fig. 15 A) shows a SEM micrograph of a wear scar induced by theball-cratering device and Fig. 15 B) the surface of a coating of a moldafter a long time in service, which allows pitting to be identified as themajor failure mechanism. The image also shows some cracks that canbe attributed to the localized melting of hard, brittle phases.

Table 4Results of the linearization analysis.

Specific wear rate,k (mm3/N m)

Average STD Confidenceinterval 90%

r2

Specimen 1 6.9×10−4 5.2×10−5 5.7×10−4 to 8.0×10−4 0.977Specimen 2 9.6×10−4 6.6×10−5 8.2×10−4 to 1.1×10−3 0.982Specimen 4 8.0×10−4 4.6×10−5 7.0×10−4 to 9.0×10−4 0.987

image of Fig.�13

image of Fig.�14

Fig. 16. SEM micrograph of: A) 3-body abrasion or rolling wear mechanism, and B) 2-body abrasion or grooving wear mechanism.

Fig. 15. SEM micrograph of: A) wear scars of specimen 1, and B) of the surface of acoating after testing in service.


Ball cratering is a micro-abrasion technique in which testconditions can be tailored to adjust the wear mechanism to grooving,which is plastic deformation controlled, or rolling, which is a fracturecontrolled process [19]. Prior tests were used to select test parameterscompatible with rolling wear, designed to induce pitting failure.

The wear scar displays two modes of wear: pitting and grooving,pitting being the more extensive. Grooving occurs when particles sliponto the surface. Fig. 16 shows the morphology of these two wearmodes in higher magnification. The wear mechanism present in thespherical cup depressions (essentially 3-body rolling wear) accuratelyrepresents the wear mechanism in the molds. However one shouldkeep in mind that the molds in service are subject to other damagemechanisms such as thermal and mechanical fatigue and corrosion athigh temperature.

4. Conclusions

The influence of the arc current used in the plasma transferred arcprocess (PTA), as well as the effect of post-weld heat treatment, on themicrostructure, hardness and wear resistance of deposited layers wasanalyzed in this study. The increase in the arc current increased thedilution of the cast iron, changing the composition andmicrostructureof the coatings. Moreover, the increase in current raises the quantityof precipitates and C flakes at the grain boundaries, which leads to therefinement of the microstructure. Increasing the arc current gives riseto a reduction in the hardness in all regions of the coatings. The

reduction in the hardness of the coatings is accompanied by areduction in their wear resistance in wear tests performed at roomtemperature with silica slurry. Post-weld heat treatment reduces thehardness in the partially melted zone and heat affected zone asopposed to the coatings themselves, where hardness remainsunaffected.

In the future, abrasion and oxidation tests at high temperature willbe done, in order to reproduce the conditions of industry application.

Acknowledgements

The authors wish to express their sincere thanks to the companyIntermolde, to Instituto Pedro Nunes (IPN) and the PortugueseFoundation for the Science and Technology (FCT) through COMPETEprogram from QREN and to FEDER for financial support.

References

[1] M. Cingi, F. Arisoy, G. Basman, K. Sesen, Mater. Lett. 55 (2002) 360.[2] E.B. Hinshaw, ASM HANDBOOK, 1993, p. 740.[3] A. Gatto, E. Bassoli, M. Fornari, Surf. Coat. Technol. 187 (2004) 265.[4] K. Siva, N. Murugan, R. Logesh, Int. J. Adv. Manuf. Technol. 41 (2009) 24.[5] V. Balasubramanian, A.K. Lakshminarayanan, R. Varahamoorthy, S. Babu, J. Iron.

Steel Res. Int. 16 (2009) 44.[6] E.H. Takano, D. Queiroz, A.S.C.M. D' Oliveira, Soldagem Insp. São Paulo 13 (2008)

210.[7] V.V. Díaz, J.C. Dutra, A.S.C. D' Oliveira, Soldagem Insp. São Paulo 15 (2010) 41.[8] K. Gurumoorthy, M. Kamaraj, K.P. Rao, A.S. Rao, S. Venugopal, Mater. Sci. Eng., A

456 (2007) 11.[9] E.O. Ezugwu, Z.M. Wang, A.R. Machado, J. Mater. Process. Technol. 86 (1998) 1.

image of Fig.�16

image of Fig.�15


[10] A.J. Ramirez, J.C. Lippold, Mater. Sci. Eng., A 380 (2004) 245.[11] A. Ramalho, Wear 269 (2010) 213.[12] M. Pouranvari, Mater. Des. 31 (2010) 3253.[13] F.A.L. Dullien, Porousmedia, Fluid transport and pore structure, Academic Press, 1992.[14] J.-F. Charmeux, R. Minev, S. Dimov, E. Minev, Borovetz, Proceedings of

International Conference, 4M2007, Whittles Publishing (2007) 217–220.

[15] Commersald, in: C.P. Technology (Ed.) Commersald PTA Technology, Modena(Italy), http://www.commersald.com, (2009).

[16] M.G. Collins, J.C. Lippold, Welding J. 82 (2003) 288S.[17] A.J. Ramirez, J.W. Sowards, J.C. Lippold, J. Mater. Process. Technol. 179 (2006) 212.[18] R.M. Imayev, V.M. Imayev, M. Oehring, F. Appel, Intermetallics 15 (2007) 451.[19] R.I. Trezona, D.N. Allsopp, I.M. Hutchings, Wear 225–229 (1999) 205.

http://www.commersald.com

Annex B


Annex B

F. Fernandes, A. Cavaleiro, A. Loureiro, Oxidation behavior of Ni-based coatings

deposited by PTA on grey cast iron, Surface and Coatings Technology, 207 (2012)

196-203.

Surface & Coatings Technology 207 (2012) 196–203

Contents lists available at SciVerse ScienceDirect

Surface & Coatings Technology

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Oxidation behavior of Ni-based coatings deposited by PTA on gray cast iron

F. Fernandes ⁎, A. Cavaleiro, A. LoureiroCEMUC - Department of Mechanical Engineering, University of Coimbra, Rua Luís Reis Santos, 3030‐788 Coimbra, Portugal

⁎ Corresponding author. Tel.: +351 239 790 700; faxE-mail address: [email protected] (F. Ferna

0257-8972/$ – see front matter © 2012 Elsevier B.V. Alldoi:10.1016/j.surfcoat.2012.06.070

a b s t r a c t
Article history:Received 28 February 2012Accepted in revised form 20 June 2012Available online 29 June 2012

Keywords:High temperature oxidationOxide scalesSurface morphologyNickel alloysPlasma transferred arcThermogravimetric measurements

The aim of this investigation was to study the effect of PTA current (100 and 128 A) on the oxidation behaviorof nickel-based hardfacing coatings deposited on gray cast iron. The oxidation behavior of coatings held at800 and 900 °C for 2 h in air was studied by thermogravimetry (TGA). The surface, aswell as the cross‐section,of the coatingswas characterized by scanning electronmicroscopy combinedwith energy-dispersive X-ray spec-troscopy (SEM/EDS) and X-ray diffraction (XRD). TGA results indicate that the coating produced with lower arccurrent exhibits more effective oxidation resistance than that producedwith higher current. This behavior couldbe correlated with the dilution promoted by the PTA process, which changes the chemical composition of coat-ings. As a consequence, different kinds of oxide scaleswere detected in each coating after isothermal oxidation. Inthe specimen producedwith a lower arc current and lower dilution a protective layer of Si–O is formed, while inthe specimen producedwith a higher current two layers could be identified: an external one of Fe2O3, with smallfeatures of Fe3O4 and NiFe2O4 spinel rich phases, and an internal one of Fe3O4with small amount of a dark phaseevenly distributed rich in silicon, nickel and iron. The isothermal oxidation curve at 900 °C of the coating withhigher dilution showed two stages: at an early stage, the weight increase over time is almost linear whereas,in a second stage, a parabolic law could be fitted to the experimental data. The other specimens followed onlya parabolic law.


1. Introduction

Cast irons and copper-alloys are commonly used in the production ofglassmolds and accessories for the glass industry [1]. In service,molds aresubjected to very severe abrasion, wear, oxidation, and fatigue at hightemperatures, due to contact with melted glass, which limit their surfacelife. As a rule, automated manufacturing processes are operated withmold temperatures close to the critical temperature of themoldmaterial.Although the temperature of melted glass is about 1050 °C the surfacetemperature of molds is currently in the range of 500–900 °C, due to aninternal cooling system in the molds. Above a critical temperature theglass body starts to adhere to the mold surface. Higher temperatureslead to mold failure. Therefore, protective coatings are normally appliedto give themolds high temperature resistance, with the aim of increasingtheir lifetime in harsh environments [2,3].

Nickel-based superalloys gain an extremely important role in protec-tion of surfaces in components for the glass industry due to their excel-lent performance under conditions of abrasion and corrosion at elevatedtemperatures [4–6]. Several processes have been used to apply coatingsto protect the surfaces of these components, as follows: (i) thermalspraying processes such as high velocity oxy fuel (HVOF), atmosphericplasma spraying (APS) and flame spraying (FS); and (ii) hardfacing pro-cesses such as plasma transferred arc (PTA) and gas tungsten arc

: +351 239 790 701.ndes).

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welding (GTAW) processes. However, since hardfacing processes giverise to a metallurgical bonding with the substrate they have been usedpreferentially in the protection of mold surfaces as opposed to thermalspraying processes which produce only mechanical bonding of thelayer to the substrate. Among the hardfacing processes above, PTA hasbeen widely used to protect the surface of molds, especially theiredges, because they are submitted to heavy wear duty cycles. This pro-cess gives rise to thicker deposits of very high quality, providing a highdeposition rate in a single layer [7–9]. However, the chemical composi-tion and properties of the coatings, as well their quality, are stronglyinfluenced by the dilution of the substrate promoted by the PTAprocess.Low dilution provides coatings with a chemical composition similar tothe added metal powder, which is a prerequisite to getting improvedwear and oxidation resistance [10]. However, as dilution is decreasedby reducing process current the occurrence of defects such as lack of fu-sion increases,which is catastrophic in terms of the behavior ofmolds inservice. So during coating deposition, the tendency is to increase the di-lution of the base material even if oxidation resistance can be under-mined. As oxidation is one of the major failure factors at hightemperature in molds for the glass industry, it is imperative to under-stand how the different coatings provided by different process condi-tions behave under the harsh and oxidizing atmospheres.

Technical literature reports a lot of studies on the high temperatureoxidation behavior of nickel alloys. Liu et al. [11], who studied the oxi-dation behavior in air of a single-crystal Ni-base superalloy at 900 and1000 °C, reported two oxidation steps for both temperatures. The first





Table 1Nominal chemical composition (wt.%) of substrate and nickel alloy.

Base material C Mn Si P S Cr Ni Mo V Ti Fe

Gray cast iron 3.60 0.60 2.00 b0.20 b0.04 b0.20 b0.50 0.50 0.10 0.20 Balance

Hardfacing C Cr Si B Fe Al F Co Ni

Ni-alloy 0.14 2.45 2.56 0.86 1.08 1.30 0.01 0.08 Balance


one is controlled by NiO growth and the second by Al2O3 growth until acontinuous Al2O3 layer formed under the previously grownNiO layer. Liet al. [12] studied the oxidation of a NiCrAlYSi overlayerwith orwithouta diffusion barrier deposited by one-step arc ion plating. They showedthat the duplex coating system exhibits a more effective protection forthe substrate, where thin and continuous scales are adhered to theoverlayer surface, and very limited oxidation and interdiffusion attacksare detected. Zhou et al. [13] studied the oxidation behavior of pure anddoped nickel alloys with different Co contents in air at 960 °C. They ob-served that increasing the Co content increases themass oxidation gainof the coatings.

Since oxidation is one of the main drawbacks that limit the life ofglass molds, the aim of this research is to study the effect of PTA currentvariation on the oxidation behavior of coatings deposited on gray castiron using a nickel-based alloy. The oxidation behavior of the coatingswas studied by thermal gravimetric analysis (TGA). The surface andcross‐section morphologies of coatings after isothermal oxidation wereobserved and characterized by scanning electron microscopy providedwith energy dispersion spectrometry (SEM-EDS) and x-ray diffraction.Further, the high temperature oxidation mechanisms are discussed.

2. Experimental procedures

Nickel based Colmonoy 215 (fromColmonoy Company) powderwasdeposited by plasma transferred arc (PTA) onto specimen blocks of graycast iron, currently used in the production of molds for the glass indus-try. The deposits were executed using a Commersald Group ROBO 90machine using two different arc currents (100 and 128 amperes). Thegoal of using two different arc currents was to achieve coatingswith dis-similar levels of dilution with the substrate. The nominal compositionsof the cast iron and Colmonoy 215 powder are displayed in Table 1.The principal surfacing parameters employed are shown in Table 2. Be-fore coating, the blocks of cast iron were induction heated at 480 °C, inorder to reduce susceptibility to cracking during coating deposition.

After coating, specimens containing cast iron and coating were re-moved from each block for microstructural analysis. Metallographicanalysiswas done using conventional procedures. Small samples for ox-idation testswere also removed from each coating. The surfaces of thesesamples were ground using number 1000 SiC abrasive paper in order toproduce even surface preparations. Further, the specimens were ultra-sonically cleaned in alcohol.

Isothermal oxidation tests were conducted at 800 and 900 °C in airfor 2 h in a thermal gravimetric analysis (TGA) machine. The air fluxused was 50 ml/min and the heating rate up to the isothermal temper-ature was 20 °C/min. After thermal exposure the samples were coolednaturally in air to room temperature. The weight gain of samples wasevaluated at regular 2 s intervals using amicrobalance with an accuracyof 0.01 mg. After oxidation the surface morphology of specimens was

Table 2Main deposition parameters for PTA weld surfacing.

Main arccurrent (A)

Powder feedrate (rpm)

Travel speed(mm/s)

Powder feed gas flowrate (l/min)

Plarat

Coating 1 100 20 2 2 2.2Coating 2 128 20 2 2 2.2

examined by scanning electron microscopy with x-ray spectroscopy(SEM-EDS). Further, the same equipmentwas used to identify the layersforming the oxide scale from the cross‐section of samples. In this casethe specimensweremounted in epoxy resin, polished and then surfacedwith a thin layer of gold for perfect observation. X-ray diffraction (XRD)using Co Kα radiation was conducted on the oxidized surface of speci-mens. In order to ensure the reproducibility of results, three specimenswere analyzed for each coating and set of test conditions.

3. Results

3.1. Microstructure of the as-deposited coatings

Fig. 1 displays the microstructure of the coatings deposited on thesurface of gray cast iron coupons. The coatings show dense micro-structure without lack of fusion of the substrate, free of microcracksand few solidification voids. Their typical microstructure consists ofdendrites of the Ni-Fe solid solution phase aligned along the directionof heat flow. Furthermore, it displays C-flakes (dark-floret like struc-tures) randomly distributed in the matrix. The increase in arc currentgives rise to a higher dilution of the gray cast iron, changing the orig-inal chemical composition of the filler metal. Following the proceduredetailed in reference [9], dilutions of 28% and 54% were measured forthe coatings deposited using 100 and 128 A arc current, respectively.These levels of dilution can explain the increase in C-flakes with in-creasing arc current, as shown in Fig. 1. Table 3 shows the chemicalcomposition measured by energy dispersive spectrometry (SEM-EDS)from an area of 400×400 μm, from the cross‐section at the middlethickness of each coating. With this procedure it is assured that chem-ical composition values are measured from a representative volume ofmaterial in that zone, avoiding erroneous measurements that couldarise from point analyses of the heterogeneous microstructure. As canbe seen, the main differences in chemical composition are related tothe higher iron, carbon and silicon contents achieved with increasingdilution. The detailed characterization of the microstructure of eachcoating can be found in a previous publication by the authors [9].

3.2. Isothermal oxidation in air

The results of the thermo gravimetric analysis performed on the coat-ings at different isothermal temperatures (800 and 900 °C) are shown inFig. 2a and b for the coatings deposited using 100 and 128 A arc currents,respectively. From now on and throughout the text, the coatings depos-ited using 100 and 128 amperes will be identified as “C100” and “C128”,respectively. The C100 coating exhibits parabolic oxidation weight gainas a function of time for both testing temperatures as does the C128 coat-ing at 800 °C. On the other hand, oxidation of the C128 coating at 900 °Cstarts with a linear increase in mass gain but after 1400 s it starts to

sma gas flowe (l/min)

Shielding gas flowrate (l/min)

Torch workdistance (mm)

Oscillation(mm)

Preheattemperature,(°C)

20 13 4 48020 13 4 480

Fig. 1. Typical microstructure of: A) coating deposited using 100 A of arc current, B) coatingdeposited using 128 A of arc current.


follow a parabolic path. The figures also reveals that for both oxidationtemperatures, coating C100 displays less oxidation weight gain thancoating C128. The dissimilar levels of oxidation observed in the speci-mens can be correlated with the dilution promoted by the PTA process,that changed their chemical composition, as shown in Table 3, whichwill interfere with the oxidation process. As mentioned before, themain differences in chemical composition of coatings are related to sili-con, carbon and, particularly, iron content. Wallwork [14] and Ezugwu[15] reported that if iron content is increased in nickel-based alloys, ittends to decrease their oxidation resistance, as is the case in our study.This behavior was attributed to a progressively higher cation diffusionrate in the scale. Therefore, increasing the dilution of gray cast iron by in-creasing the arc current proves to be detrimental, since the oxidation re-sistance of coatings is reduced due to a further increase of iron content inthe coatings. Moreover, it is observed that the oxidation gain of speci-mens increases with increasing isothermal temperature. This is anexpected effect since all diffusion phenomena ruling the oxidation pro-cess are enhanced as the temperature is increased. In conclusion, thethermo gravimetric results show that increasing dilution decreases theoxidation resistance of the coatings.

3.3. X-ray diffraction

Figs. 3 and 4 show the XRD spectra found at the oxidized surface ofthe coatings. The oxide products revealed differences between the twocoatings; however, in each coating the phases produced after oxidationat 800 and 900 °C are similar, although different oxidation weight gainswere measured. According to XRD patterns, the main oxide phase sets

Table 3Nominal chemical composition (wt.%) analyzed at middle thickness of coatings depos-ited using 100 and 128 A arc current.

Elements Ni Fe C Si Al Cr

Coating deposited using 100 A 85.40 8.47 0.30 2.17 0.81 2.86Coating deposited using 128 A 51.09 43.29 1.01 2.54 0.50 1.87

detected at 800 °C and 900 °Cwere: (i) B2O3, SiO2 andOAlB for coatingsproduced with lower dilution and (ii) Fe2O3 (hematite), Fe3O4 (magne-tite) and spinels of NiFe2O4 for higher dilution coatings. The indexationof oxide phases of coating C128 is in agreement with the results fromthe literature, where similar oxide phases were detected after thermalexposure [16,17]. Both coatings also display a high intensity peak closeto 60°, which corresponds to the face centered cubic (Ni-Fe) solid solu-tion (ICDD card 47‐1405). Furthermore, coating C100 displays low in-tensity XRD peaks which correspond to a Ni–Si phase, such as Ni3Si(ICDD card 32‐0699), and a Cr-boride phase Cr5B3 (ICDD card 32‐0278). All these phases are currently identified in this type of micro-structure [9]. The large amount of iron oxide detected at the oxidizedsurface of coating C128 can be attributed to the higher content of ironin this sample. As referred to before, the higher iron content in this coat-ing is attributed to higher dilution of the base material induced by theuse of a higher arc current. Comparing the peaks intensity of coatingC128, oxidized at 800 °C and 900 °C, the amount of oxide productsFe3O4, Fe2O3 and NiFe2O4 increased with temperature. The NiFe2O4 spi-nel phase (ICDD card 10‐0325), identified at the oxidized surface, resultsfrom the reaction of NiO (whichwas also identified as an oxide phase onthe surface of coating C128) with the Fe2O3 phase, as reported byMusićet al. [18]. Furthermore, B-O phases (ICDD card 06‐0297 (B2O3) andICDD card 50‐1505 (B6O)) could be identified from the XRD pattern ofcoating C128.

3.4. Surface morphology

The different chemical compositions of coatings brought about by di-lution led to the formation of different kinds of oxide scales during iso-thermal oxidation, as shown above. The SEM surface morphologies ofthe oxidized coatings at 900 °C are shown in Figs. 5 and 6 for coatingsC100 and C128, respectively. EDS examinations were conducted of theoxidized surface of coatings to characterize their composition. CoatingC100, oxidized at 900 °C, displays two phases: a gray dark phase, identi-fied in Fig. 5a by letter A, and an evenly distributed white phase,indentified by letter B. A magnification of each zone is shown in Fig. 5b

Fig. 2. Isothermal oxidation curves at 800 and 900°C of: a) coating deposited using 100A arc current, b) coating deposited using 128 A arc current.

image of Fig.�2

Fig. 4. X-ray diffraction patterns of oxidized surface of coating deposited using 128 Aarc current, after oxidation at 800 and 900°C.

Fig. 3. X-ray diffraction patterns of oxidized surface of coating deposited using 100 Aarc current, after oxidation at 800 and 900°C.


and c in order to better illustrate the phases. EDS analysis reveals that thegray phase is rich in boron and oxygen, suggesting that it is a boronoxide, as identified by XRD analysis. A similar phase was detected by Liand Qiu [19], and Lavrenko and Gogotsi [20] after oxidation of boroncarbide powder in air, ranging from 500 to 800 °C and, oxidation ofhot-pressed boron carbide in air up to 1500 °C, respectively. They ob-served that even if the isothermal oxidantion temperature exceeds themelting point of the boron trioxide (about 577 °C), after cooling it is pos-sible detect cristaline boron oxide. According to them, at a temperatureclose to 577 °C the boron oxide melts, covering the oxide surface andserving as protective layer. Therefore, oxygen is inhibited from passinginward and metal atoms from passing outward through the oxidelayer. Above a threshold temperature, liquid boron oxide starts to evap-orate, but only after1000 °C does this evaporation becomes significant.Therefore, when the oxidation temperature is below 1000 °C, duringcooling down to room temperature, it is expected that boron oxidewill solidify on top of the oxidized surface of the material. The whitephase appears associated with spalling zones and cracks. The EDS anal-ysis shows that different phases coexist in this zone. A zone rich inboron (see peak ID 2) and another rich in silicon (see peak ID 3) wereidentified, as shown in Fig. 5c. These results matchedwith XRD analysiswhich suggest that the main phases of the oxidized surface are boronand silicon oxides, which are responsible for oxidation resistance ofthe coating due to their protective character [17,19–21]. Moreover, afloret-like structure appears to be frequently associated with thewhite phase, as shown in Fig. 5d. The EDS results (see peak ID 4)show that the elemental composition of this oxide is rich in aluminium.Taking into account the XRD results it can be concluded that the mainphase of the floret-like oxide is OAlB. The morphology and indexationof this phase named aluminum borate whiskers, matches well withthat studied in references [22,23].

However, analysis of the oxidized surface of coating C128 at 900 °Creveals that the oxide morphology is different from that of coatingC100 (see Fig. 6a). It displays amicrostructurewith uniformly distribut-ed irregular ditches and dissimilar phases. Amagnification of themicro-structure is shown in Fig. 6b to make this clearer and more easilycharacterized. The EDS analysis reveals that the phase with irregularditches (peak ID 1) is rich in iron and oxygen, suggesting an ironoxide, as identified in the XRD pattern. Further XRD analysis revealedthat the oxidized surface is composed primarily of Fe2O3 with smallamounts of Fe3O4. This result is in agreement with that of Guo et al.[16]. It is possible to identify other small features such as anotherphase rich in iron (see peak ID 2 relating to the angular particles) forwhich EDS analysis shows the presence of nickel. The combination ofNi with the Fe2O3 phase can give rise to NiFe2O4 spinels which werealso identified by XRD. Furthermore, particles of carbon (dark phase)

could be detected at the oxidized surface of the coating, as shown bypeak ID 3. Finally a phase rich in nickel, oxygen, carbon, silicon andboron (see peak ID 4) was detected at the oxidized surface of coating128 A explaining the Ni–O, and B–O phases identified in the XRDpattern of the specimen.

3.5. Cross‐section of oxidized samples

The microstructures of cross-sections of the oxide scale of coatingsC100 and C128, after oxidation and testing at 900 °C for 2 h, are pres-ented in Figs. 7a and b, respectively. The oxide scale is about 5 μmthick for the specimen coated using the lower arc current and 10 μmthick for that using the higher current. Decarburization of coatingC128 close to the surface is also revealed. Comparing this cross‐sectionwith the oxide scale of the same coating oxidized at 800 °C (Fig. 8) it ispossible to conclude that at this temperature there is no decarburizationof the specimen. Zones of silicon segregationwere found near the inter-face in both coatings (C100 and C128) annealed at either temperature,as shown, as an example, by the EDS analysis performed on the phasemarked with number 1 in Fig. 7c for coating C128 oxidized at 900 °C.

The structure of the oxide scale will be discussed separately for eachcoating, in order tomore clearly illustrate its composition and the role ofthermodynamic stability and transport mechanisms in the establish-ment of layered structures.

Fig. 9 shows SEM images in secondary and backscattered electrons ofcross‐sections of the oxide scales of coatings C100 and C128 at 900 °C.The oxide layers are complex and composed of a mixture of phases.EDS analyses were conducted on the cross‐section of the oxide scales,in order to characterize the different oxides formed during isothermaloxidation. Table 4 summarizes the EDS analyses performed at the pointsmarked in the figures. The values plotted in this table indicate the ratiobetween the peak intensity of each element with the peak intensity ofoxygen for each EDS spectrum. This means that at each point analyzed,a relative increase in this ratio for one particular element is a result ofthe preferential formation of the oxide of that element over the others.Combining SEM imageswith EDS results, plotted in Table 4 it can be con-cluded that the oxide scale of coating C100 ismainly composed of siliconoxide (as the phase identified with point 1 suggests), with smallamounts of a dark phase rich in nickel, silicon, aluminum and boron(see phase identified with point 2) uniformly distributed through thelayer. This agrees well with the XRD pattern and the indexed Si-O,Ni-Si and O-Al-B phases detected. In addition, it is possible to identify aphase rich in boron, normally located at the top of the oxidized surface(point 3), suggesting a boron oxide, as also revealed by XRD. Moreover,a Ni-rich white phase (see point 4) appears at the interface betweenthe coating and the oxide layer, corresponding to the remaining part of

image of Fig.�4

image of Fig.�3


the original source of elements which diffused outwards to form oxidescales.

The oxide scale of coating C128 is different from that of coating C100,being composed of two layers. Combining the XRD analysis, with theSEM-EDS results, the external layer can be identified as Fe2O3, the mainphase containing small features of Fe3O4 and spinel NiFe2O4. On theother hand, the internal layer is composed of a continuous Fe–O richphase (phase identified with point 2) with small evenly distributedamounts of a dark oxide rich in silicon, nickel and iron, as point 3 ofEDS analysis shows. The lower iron content in the internal layer in rela-tion to the external one suggests that the internal Fe–O rich phase re-spects the Fe3O4 oxide. This result is coherent with the oxidation ofFe-based materials and so hematite (Fe2O3) and spinel of NiFe2O4 canbe expected to form in the region of higher oxygen potential while mag-netite (Fe3O4) will form in the lower oxygen potential region [24]. Simi-lar iron oxide layers were observed by Guo et al., after oxidation of Fe–36% Ni bicrystals in air at high temperature [16]. Between these outerand inner layers, a continuous thin layer of a dark phase appears at theinterface, as revealed by SEM. A boron rich phase could be detected atthe top of the outer layer (see point 4), which is in good agreementwith the boron oxide detected by XRD. Finally, clusters of carbon are dis-tributed evenly in the oxide layers. The accumulation of C at the interface

Fig. 5. A) SEM observation of surface morphology of the coating deposited using 100 A arc ctified in A), D) magnification of the zone C identified in C). SEM/EDS spectra of: E) zone 1,

between the coating and the oxide scale (point 5) is a consequence of theoxide layer serving as a diffusion barrier for C coming from the decarburi-zation of the sample mentioned above. Analogous oxide products wereobserved in the cross‐sections of the oxide scale of coating C128 oxidizedat 800 °C.

A good correlation was found between the scale microstructure andthe isothermal oxidation curve of coating C100 oxidized at 900 °C. Theoxidation kinetics of metals and alloys is often controlled by a parabolicrate law at high temperature [25] due to the presence of protective scales.This may well be the case in this specimen, as the evolution of the oxida-tionweight gain is controlled by the silicon oxide, themain phase presentin the oxide scale, as suggested by Fig. 9. At lower temperatures, a boronoxide starts forming on the surface of the specimen due to the high affin-ity of boron to oxygen. However, for higher temperatures, due to the in-tense segregation of silicon through the interface and its great affinity foroxygen, a silicon oxide layer starts to be formed beneath the boron oxide.At the same time small amounts of Ni–Si and O–Al–B phases are formeddue to the diffusion of other elements through the interface. At a temper-ature of approximately 577 °C, boron oxide melts, covering the oxidizedsurface of the coating and inhibiting the transport of species though themelted oxide layer. Above this temperature it is expected that meltedB2O3 starts to evaporate: however, this vaporization does not

urrent, after 2 h oxidation at 900 °C, B) and C) magnification of the zones A and B iden-F) point 2, G) point 3, H) zone 4.

image of Fig.�5

Fig. 7. Back-scattered SEM image of the cross‐section of the oxide scale obtained at 900°C of: A) coating deposited using 100 A arc current, B) coating deposited using 128 A arccurrent. C) Segregation of silicon in the coating deposited using higher arc current, oxidized at 900°C. D) SEM EDS spectra of point 1.

Fig. 6. A) SEM observation of surface morphology of the coating deposited using 128 A arc current, after 2 h oxidation at 900°C, B) magnification of the microstructure in A). SEM EDSspectra of: D) point 1, D) point 2, E) point 3, F) point 4.


image of Fig.�7

image of Fig.�6

Table 4Ratio between the peak intensity of each element with the peak intensity of oxygen (O Ka).

Point B Ka C Ka Fe La Ni La Al Ka Si Ka

Coating C100 1 0.12 0.14 – 0.16 0.34 0.602 0.10 0.15 0.23 0.58 0.34 0.633 1.61 – – – 0.25 0.894 0.44 0.83 2.57 6.06 – 0.94

Coating C128 1 – 0.06 0.33 – – –

2 – 0.04 0.25 0.16 0.05 0.053 – 0.11 0.23 0.13 0.10 0.314 1.00 1.33 0.31 – – 0.375 – 1.84 0.16 0.23 0.08 0.20

Fig. 8. SEM images of the cross‐section of the oxide scale obtained at 800°C of thecoating deposited using 128 A arc current: A) SEM secondary electron (SE) image,B) back-scattered SEM image.

Fig. 9. SEM images of the cross‐section of the oxide scale of the coatings depositedusing100 and1secondary electron (SE) images.


significantly influence the slope of the kinetics curve at 800 and 900 °C,since it only has an important impact at temperatures above 1000 °C[19,20]. Finally, due to the continuous segregation of silicon, its oxidelayer will progressively thicken. A protective layer was thereforeestablished and the reaction rate becomes governed by the rate of speciesdiffusion through this thicker silicon oxide layer, as observed in the cross‐section of the oxide scale of the specimen. As with coating C100, a strongcorrelation can be found between the microstructure of the oxide scaleand the isothermal oxidation curve of coating C128. At 800 °C, parabolicbehavior is displayed; however, at 900 °C, two steps can be detected, seeFig. 2b. The isothermal oxidation curve at first exhibits a linear increase inmass gain and then starts to conform to a parabolic law. The microstruc-tural analysis of the cross-section of the oxide scale revealed that at800 °C there is no decarburization of the coating, which is not whatwas observed at 900 °C. Like the specimen coated using the lower arccurrent, at low temperatures boron oxide starts forming on the surfaceof specimen due to the high affinity of boron to oxygen. Moreover, forhigh temperatures, due to the high iron content, a consistent layer ofFe2O3 with small features of Fe3O4 and NiFe2O4 is formed, which is thick-ened by outward Fe diffusion. At the same time the specimen loses car-bon by decarburization, due to the formation of CO2. However, carbon

28A arc current oxidized at 900°C: A) andC) in back-scattered SEM images, B) andD) in SEM

image of Fig.�8

image of Fig.�9


has increasing difficulties diffusing outwards through the oxide layer andstays either inside or beneath the oxide scale, as shown bymorphologicalanalysis of the cross-section. Therefore, after the first step where theweight gain is partially compensated by C liberation, the oxidationcurve will deviate to a normal parabolic trend. According to theRichardson-Ellingham diagram, Si has higher affinity for O than Ni andFe, and Fe forms Fe3O4 more easily than Ni forms NiO. Therefore, duringoxidation it would be expected that Si-O would be the first layer to beformed. However, its much lower content in the specimen relative toiron, indicates that considerable time is required to form a continuousand protective SiO2 layer, as suggested by Douglass and Armijo [26]. So,after a critical time a layer of SiO2 starts being formed below the Fe2O3

scale, but it never becomes thick enough to determine the oxidation ki-netics. Further, due to the continuous segregation of iron through the sur-face of the specimen and due to the decrease in oxygen content at thecoating's surface as a result of the volume growth of the Fe2O3 layer, acontinuous and consistent Fe3O4 layer starts forming in this zone, withsmall uniformly distributed amounts of the dark phase.

4. Conclusion

This investigation concerned the effect of PTA current variationon the oxidation behavior in air at 800 and 900 °C of coatings of anickel-based hardfacing alloy deposited on gray cast iron using two dif-ferent arc currents (100 and 128 A). Samples were analyzed by thermalgravimetric analysis, x-ray diffraction (XRD) and scanning electron mi-croscopywith energy dispersion spectrometry analysis (SEM-EDS). Thethermo gravimetric results show that increasing the arc current from100 to 128 A decreases the oxidation resistance of coatings, a resultthat can be correlated with the different chemical compositions of thecoatings, promoted by different dilutions of the base material. Thecoating with lower dilution follows a parabolic oxidation weight gainas a function of time, as this behaviour is essentially controlled by thegrowth of a Si–O layer. The coating with higher dilution follows thesame trend at 800 °C. However, at 900 °C two stages in the oxidationcurvewere detected, the firstwith a linear increase inmass gain, a com-promise between outward Fe diffusion through a growing layer ofFe2O3 and the loss of carbon by decarburization and formation of CO2.The second step obeys a parabolic law from the moment that C libera-tion is impeded by a thickened scale consisting of an external Fe2O3

phase, with small amounts of Fe3O4 and NiFe2O4 spinel rich film, and

an internal layer of Fe3O4 film with small evenly distributed amountsof a dark phase rich in silicon, nickel and iron. In summary, an increasein arc current from 100 to 128 A in the deposition of nickel-rich layerson substrates of cast iron by the PTA process is very harmful, as itreduces the oxidation resistance of the coatings at high temperature.

Acknowledgements

The authors wish to express their sincere thanks to the PortugueseFoundation for the Science and Technology (FCT), through COMPETEprogram from QREN and to FEDER, for financial support in the aim ofthe project number “13545”, as well as for the grant SFRH/BD/68740/2010.

References

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456 (2007) 11.[6] W. Li, Y. Li, C. Sun, Z. Hu, T. Liang, W. Lai, J. Alloys Compd. 506 (2010) 77.[7] A. Gatto, E. Bassoli, M. Fornari, Surf. Coat. Technol. 187 (2004) 265.[8] K. Siva, N. Murugan, R. Logesh, Int. J. Adv. Manuf. Technol. 41 (2009) 24.[9] F. Fernandes, B. Lopes, A. Cavaleiro, A. Ramalho, A. Loureiro, Surf. Coat. Technol.

205 (2011) 4094.[10] V. Balasubramanian, A.K. Lakshminarayanan, R. Varahamoorthy, S. Babu, J. Iron

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Res. 8 (2005) 365.[25] Y. Nakamura, Metall. Mater. Trans. A 6 (1975) 2217.[26] D.L. Douglass, J.S. Armijo, Oxid. Met. 2 (1970) 207.

Annex C


Annex C

F. Fernandes, A. Ramalho, A. Loureiro, A. Cavaleiro, Wear resistance of a nickel-

based coating deposited by PTA on grey cast iron, International Journal of Surface

Science and Engineering, 6 (2012) 201-213.

Int. J. Surface Science and Engineering, Vol. 6, No. 3, 2012 201

Copyright © 2012 Inderscience Enterprises Ltd.

Wear resistance of a nickel-based coating deposited by PTA on grey cast iron

Filipe Fernandes* CEMUC – Department of Mechanical Engineering, University of Coimbra, Rua Luís Reis Santos, 3030-788 Coimbra, Portugal and Intermolde, Rua de Leiria, 95, Apartado 103, 2431-902 Marinha Grande, Portugal Fax: +(351) 239-790-701 E-mail: [email protected] *Corresponding author

Amilcar Ramalho, Altino Loureiro and Albano Cavaleiro CEMUC – Department of Mechanical Engineering, University of Coimbra, Rua Luís Reis Santos, 3030-788 Coimbra, Portugal Fax: +(351)-239-790-701 E-mail: [email protected] E-mail: [email protected] E-mail: [email protected]

Abstract: The moulds for the production of glass bottles, made of cast iron, are subjected to very severe conditions of wear during use. Thus, it is essential to understand the wear mechanisms involved, in order to increase the equipment life. The present study intends to study the abrasion resistance of moulds submitted for long time at working conditions. The investigation was conducted on nickel-based coatings deposited by plasma transferred arc (PTA) on grey cast iron, using different arc currents. Micro-scale ball cratering abrasive wear test was used to evaluate the tribological properties of the as-deposited and heat treated coatings. The results show that micro-scale abrasion tests induce grooving and pitting wear mechanisms. Increasing the arc current decreased the hardness of the coatings and, consequently, their wear resistance. The hardness and wear resistance of the coatings was improved by increasing heat treatment holding time.

Keywords: plasma transferred arc; PTA; Ni-based alloy; wear; ageing; glass moulding; surface engineering.

Reference to this paper should be made as follows: Fernandes, F., Ramalho, A., Loureiro, A. and Cavaleiro, A. (2012) ‘Wear resistance of a nickel-based coating deposited by PTA on grey cast iron’, Int. J. Surface Science and Engineering, Vol. 6, No. 3, pp.201–213.

202 F. Fernandes et al.

Biographical notes: Filipe Fernandes is a PhD student in the Mechanical Engineering Department of the University of Coimbra. He obtained his MSc in 2009. He is currently dedicated to the surface engineering field applied to the moulds for the glass industry. He is focused on the study of new methods to protect the surface of those components. His research interests include thermal spraying of coatings (PTA, HVOF, APS and PVD), oxidation and wear resistance of materials suitable for high temperature applications.

Amilcar Ramalho received his PhD in Mechanical Engineering from Universidade de Coimbra, Portugal in 1994. He is currently a Professor of Maintenance and Mechanical Vibrations. His research interests include friction and wear of materials and components, fretting, tribotesting and biotribology. He is the co-author of more than 150 papers in international journals and conferences. He belongs to the editorial board and collaborates as a reviewer of several international journals.

Altino Loureiro is currently Associate Professor of Welding and Related Technologies in the University of Coimbra (UC) and Researcher from the Center of Mechanical Engineering at the same university (CEMUC). His research interests are focused on technological and metallurgical aspects of welding and thermal spaying processes. He is the author or co-authored over 150 articles in international journals and conferences and reviewer of several international scientific and technical journals. His is also coordinated and collaborated on more than a dozen of research projects and is a consultant to companies in the field of metal construction.

Albano Cavaleiro is full Professor in the Mechanical Engineering Department of the University of Coimbra. He is dedicated to the surface engineering field for more than 25 years. He is mainly concerned with the deposition and characterisation of sputtered coatings for mechanical applications, such as, low friction and self-lubricant, high temperature resistant, oxidation and corrosion protective and DLC. He published more than 150 papers in international peer reviewed journals and presented more than 30 invited talks in international scientific events.

1 Introduction

Cast iron is commonly used in the production of moulds and accessories for glass industry. Wear and oxidation are the main failure mechanisms that limit the surface life of moulds (Cingi et al., 2002). Therefore, it is necessary to protect mould surfaces with hard and wear resistant materials.

The plasma transferred arc (PTA) process has been usually applied to coat the surfaces of moulds and, especially, their edges where the wear is higher. It produces very high quality deposits, offering optimal protection with minimal thermal distortion of the parts, and provides high deposition rate (Gatto et al., 2004). It allows very precise deposit layers of complex alloys on mechanical parts that are subjected to harsh environments, significantly extending their service life (Siva et al., 2009). However, the hardness and the wear resistance of the coatings are strongly influenced by the dilution of the substrate promoted by the PTA process. Low dilution provides coatings with similar chemical composition to the metal powder added, condition to have an improved wear and

Wear resistance of a nickel-based coating deposited by PTA on grey cast iron 203

corrosion resistance (Balasubramanian et al., 2009). Therefore, it is essential to control the relevant process parameters to minimise the dilution.

A significant number of coating materials, such as cobalt, iron and nickel alloys, can be applied by PTA process to extend the working life of tribological components subjected to wear. In glass industry nickel alloys gain extremely important role in protection of surfaces, due to their excellent performance under conditions of abrasion and corrosion at elevated temperature. The microstructure and wear behaviour of nickel-based hard-facing alloy deposits have been studied using various alloy compositions and different substrates. Nickel-based alloys are usually strengthened by alloying elements which improve their wear properties (Ezugwu et al., 1998; Chang et al., 2010). Technical literature reports a lot of tribological studies done by pin-on-disc tests in deposits of Ni alloys on stainless and carbon steels, though the sliding contact conditions of pin-on-disc do not reproduce the interaction conditions between melted glass and the surface of the mould. Because of that, the ball cratering wear test was selected as the first step to investigate the micro-abrasion resistance of the coatings at room temperature.

Guo et al. (2011) studied the high temperature wear resistance and wear mechanism of NiCrBSi and NiCrBSi/WC-Ni coatings deposited on stainless steel by laser cladding. They found that the wear resistance of the coatings is much greater than that of stainless steel when sliding against Si3N4 counterpart ball. Furthermore, NiCrBSi/WC-Ni composite coating showed better wear resistance at high temperature than NiCrBSi coating. Kesavan and Kamaj (2010), who worked with a Colmomoy 5 hard-facing alloy deposited by PTA on a stainless steel, studied the room and high temperature wear behaviour using pin-on-disc wear tests. They showed that the wear resistance of the coatings improved significantly with increasing test temperature. Further, they observed different operating wear mechanisms depending on the sliding and testing temperature. In turn, Gurumoorthy et al. (2007) studied the sliding wear of a nickel hard-facing alloy AWS NiCr-B deposited on the same stainless steel at three different temperatures (room temperature, 300°C and 500°C), and the influence of an ageing treatment on its microstructure, hardness and wear behaviour. The investigation revealed a significant weight loss at room temperature and an abrupt decrease in mass loss at high temperature. However, after ageing treatment they demonstrated that there is no significant change in the weight loss or the wear behaviour at room and high temperature.

Since wear and oxidation are the main drawbacks that limit the surface life of glass moulds, the aim of this research is to study the effect of PTA current on the microstructure, hardness and three-body abrasion behaviour of a nickel-based hard-facing alloy coating deposited on grey cast iron. Ageing studies were also conducted to evaluate the influence of temperature holding time on the microstructure, hardness and tribological properties of the coating processed with the highest current.

2 Experimental procedure

A nickel-based hard-facing alloy named Colmonoy 215 (from Colmonoy Company) was deposited on flat surfaces of grey cast iron blocks by PTA process. Before coating, each cast iron block was preheated at 480°C, in order to reduce the thermal shock and cooling rate to avoid cracks in both the coating and heat-affected zone. The same process


parameters were used for all coatings except arc current, which was 100 A for specimen 1, 128 A for specimen 2 and 140 A for specimen 3. The composition of cast iron and metal powder is given in Table 1 and the process parameters used in Table 2. After coating, the blocks were left to cool to room temperature. Samples were removed from each block, transversely to welding direction, for further analysis. Ageing studies were carried out at 500°C for 5, 10 and 20 days on specimen 3. Table 1 Nominal chemical composition (wt.%) of substrate and hard-facing alloy


Grey cast iron 3.60 0.60 2.00 < 0.20 < 0.04 < 0.20 < 0.50 0.50 0.10 0.20 Balance Hard-facing C Cr Si B Fe Al F Co Ni Ni-alloy 0.14 2.45 2.56 0.86 1.08 1.30 0.01 0.08 Balance

Table 2 Deposition parameters for PTA weld surfacing.

Main arc current (A)

Powder feed rate (rpm)

Travel speed (mm/s)

Powder feed gas flow rate (l/min)

Specimen 1 100 20 2 2 Specimen 2 128 20 2 2 Specimen 3 140 20 2 2

Plasma gas flow rate (l/min)

Shielding gas flow rate (l/min)

Torch work distance

(mm)

Oscillation (mm)

Preheat temperature

°C Specimen 1 2.2 20 13 4 480 Specimen 2 2.2 20 13 4 480 Specimen 3 2.2 20 13 4 480

Microstructural analysis was done in cross section of the samples. Optical and scanning electron microscopies (OM and SEM) were used to characterise the microstructure of the coatings. Vickers micro-hardness profiles were measured in the cross section of treated samples using a load of 5 N.

Micro abrasive wear tests were carried out in as-deposited and heat treated surfaces at room temperature in a ball cratering devices. Figure 1 shows a schematic diagram of the equipment used. In this test a specific normal load is applied to press a rotating ball into the flat surface of the coated sample in the presence of a slurry suspension of abrasive. The abrasive slurry, agitated by a magnetic stirrer, was continuous and gravitationally drip fed onto the rotating ball to produce a spherical cup depression. Preliminary tests using different contact conditions were done to reproduce the interaction conditions in the surface of the moulds which have been in service for a long time. The wear mechanism was studied by SEM. High purity silica SS40 with angular particles of averaging of 3.10 μm in size was used as an abrasive to replicate the effect of the melted glass. The slurry used was prepared with a concentration of 103 g per 1 × 10–4 m3 of distilled water. A steel ball bearing (AISI 52100) with 0.0254 m was used. Several scars were made in each sample with different durations, 100, 200, 300, 400 and 500 turns, respectively 7.98, 15.96, 23.94, 31.92, 39.90 meters of sliding distance for all the specimens, under the action of a normal load of 0.1 N. In order to minimise the scatter of results, previously to starting the tests, the ball surface was etched with chloridric acid (33%) during ten


minutes. The dimensions of spherical depressions were measured in order to calculate the wear volume, which is given by the equation (1). In this equation, R represents the ball radius and b the crater chordal diameter of the spherical cup depression. The b value was evaluated using the average of two perpendicular measurements performed at the wear scar chordal diameter. Those values were measured with a Mitutoyo Toolmaker Microscope with x-y micrometer table.

( )4V π b (64 R)= × × × (1)

In order to obtain the specific wear rate, to quantify the wear behaviour, Archard’s law was applied, see equation (2). This is the most applied law to determine a parameter to quantify the wear behaviour of materials. In this equation, P is the normal load, l is the sliding distance and k is the specific wear rate.

V k P l= × × (2)

To quantify the specific wear rate a linear approach of Archard’s model was applied and a statistical analysis was used to estimate the error of measurements, as described elsewhere (Ramalho, 2010).

Figure 1 Schematic diagram of the micro-scale abrasion tests

3 Results

3.1 Microstructure

Figure 2 shows the microstructure of the coatings performed using different arc currents [Figures 2(a), 2(b) and 2(c)], as well as the fusion boundary of specimen 3 [Figure 2(d)], and SEM pictures of specimen 3 before and after heat treatment [Figures 2 (e) and 2(f)]. The typical microstructure of deposits consists of dendrites of Ni-Fe solid solution phase with columnar morphology oriented along the direction of heat flow, with C-flakes (dark floret-like structures) randomly distributed. However, near the interface it was observed the presence of a cellular microstructure finer than the dendrite structure with diminution of the size of C-flakes, as illustrated in Figure 2(d) for specimen 3. This behaviour could be explained by the higher solidification rates involved in the boundary because of the


efficient thermal exchange ensured by the high volume ratio substrate/coating of the base material, as suggested by Gatto et al. (2004) and Navas et al. (2006). Moreover, the presence of the ‘dark’ phase on the coatings cannot be explained by the chemical composition of the metal powder added, so it can be only attributed to the high dilution of cast iron induced by the PTA process. The dilution was evaluated by the ratio between the area of the melted base material and the total area of the melted zone, both measured in the cross section of the coatings. Dilutions of 28%, 50% and 59% were measured respectively for specimen 1, 2 and 3. These results agree with the observed increasing proportion of C-flakes for higher arc currents, as can be seen in Figures 2(a), 2(b) and 2(c). Therefore, the chemical composition of the coatings is largely influenced by the dilution of cast iron. A detailed characterisation of this type of coatings can be found in a previous publication of the authors (Fernandes et al., 2011).

Figure 2 Microstructure of (a) specimen 1, (b) specimen 2, (c) specimen 3, (d) interface of specimen 3 and SEM pictures of sample 3 (e) in non-heat-treated condition and (f) in heat treated condition during ten days


Table 3 SEM chemical analysis of the deposit and base material of specimen 3

Elements wt%

C Al Si Cr Fe Ni

1 0.8 0.7 2.6 1.4 54.7 39.8 Deposit 2 0.8 0.7 2.8 1.6 54.8 39.2 3 0.9 1.4 2.5 1.9 50.4 42.8 4 1.6 1.1 2.2 1.4 59.4 34.2 Interface 5 3.9 - 2.7 - 93.3 - Base material

Figure 3 SEM magnification of microstructure of the coating 3 (a) in non-heat treated condition and (b) in heat-treated condition after ten days exposed at 500°C and SEM EDAX spectra of (c) point 1, (d) point 2, and (e) point 3 (see online version for colours)

(c) (d)

(e)

Table 3 shows the chemical composition in area measured by SEM-EDS at various zones on the coating and the substrate of specimen 3. Zone 1 and 2 are located close to the coating surface, zone 3 at the mid-thickness of the coating, zone 4 in the coating close to the fusion boundary and zone 5 in the bulk cast iron. This analysis showed that PTA process promotes an evenly distribution of the elements inside the coating. The same observation was done in specimens 1 and 2. Due to the dilution, the amount of iron in the


coating increases reducing its nickel content. Ezugwu et al. (1998) reported that the increase of iron content in the coatings tends to decrease their oxidation resistance, because of the formation of a less adherent oxide scale.

Figures 2(e) and 2(f) reveals that the heat treatment induces significant changes in the original microstructure of specimen 3. It promotes the formation of a white phase on the grain boundaries. Magnifications of the two microstructures are shown in Figure 3. SEM-EDS spectra analyses were conducted on the coatings to give a detailed characterisation of their microstructure. From the EDS analysis, it can be stated that the grain boundaries of the non-heat-treated sample is formed by two phases: a light grey and a dark grey rich in chromium and silicon, respectively. Inversely, the specimen in heat-treated condition is formed by the light grey phase as observed in the non-heat-treated sample and a new phase (white phase) which is replacing the dark phase. The amount of the white phase grows with increasing time of heat treatment. According to Fernandes et al. (2011), who studied a nickel alloy deposited on grey cast iron, the dark grey phase is a Ni3Si phase. The EDS spectrum reveals that both phases (the dark and the white) are rich in Ni-Si. However, the white phase presents a higher content of phosphorous on the heat treated sample suggesting a dispersion of Ni-P precipitates on the grain boundaries, as observed by Sahoo and Das (2011) in electroless nickel coatings.

Figure 4 Hardness profile across the interface of (a) as-deposited specimens and (b) specimen 3 before and after heat treatment (see online version for colours)

(a) (b)

3.2 Hardness

The micro-hardness profiles of the different as-deposited and heat treated specimens are shown respectively in Figures 4(a) and 4(b). Figure 4(a) reveals that the hardness through the coatings is approximately constant and tends to decrease with increasing arc current. The dilution induced by the PTA process can explain this behaviour. The partial melted zone (PMZ) displays the highest hardness. According to Pouranvari (2010), who studied the weldability of grey cast iron using nickel-based filler metals, the highest hardness observed in PMZ is due to the formation of hard and brittle phases during surfacing, such as martensite and white cast iron. The heat treatment promotes the progressive hardening of the coatings with increasing thermal exposure time [see Figure 4(b)]. Some scatter of the hardness values can be observed too. Harsha et al. (2008), who studied the influence of CrC in a nickel flame sprayed coatings, observed similar behaviour. According to them, the scatter in the values is due to the presence of a harder eutectic network around the cells. As described before, heat treatment promotes the formation of Ni-P precipitates


on the grain boundaries. The greater hardness observed on the ageing treated samples can be attributed to the presence of these precipitates similarly to what is currently observed in Ni-P electroless coating when annealed at increasing temperatures (Sahoo and Das, 2011). This type of coatings is a supersaturated alloy in the as-deposited state and can be strengthened by the precipitation of nickel phosphide crystallites. The phosphides act as barriers for dislocation movement and improve the mechanical strength of the coatings.

3.3 Wear behaviour of coatings

Micro-scale abrasion tests were performed on the coatings at room temperature to predict the effect of the different process parameters on their wear behaviour. The goal in a first step was to study the abrasion resistance present in the surface of in-service moulds, in order to quantify the wear loss of coatings. The response of materials under micro abrasion tests depends on the nature of the motion of particles in the contact zone, which determines the stress state arising near the contact. Three-body wear mechanism (rolling or pitting) and two-body wear mechanism (grooving), or a combination of both mechanisms could be detected. Rolling occurs when the abrasive particles are able to free rotate between the sphere and the material in test; on the other hand, grooving occurs if the particles follow the rotation of the sphere, scratching the specimen surface (Ramalho et al., 2011; Gant and Gee, 2011). In spite of this general tendency to the occurrence of three-body or two-body abrasion, the characteristics of the materials will influence the process; mild materials will stabilise the grooving mode.

Figure 5 Evolution of the wear volume as function of the parameter ‘normal load × sliding distance’ of (a) as-deposited coatings and (b) specimen 3 before and after thermo exposure at 500°C during several days (see online version for colours)

(a) (b)

The results of the micro-scale abrasion tests from specimens in as-deposited and heat treaded conditions are plotted respectively in Figures 5(a) and 5(b), as the material volume loss as a function of the product of the sliding distance (l) by the normal load (P). The reliability analysis of the data is presented in Table 4 and the confidence interval for the specific wear rate plotted in Figure 6. For each sample, a straight linear trend was fitted to the experimental points in order to obtain the specific wear rate, which corresponds to its slope [see equation (2)]. From Figure 5(a), it is possible to conclude that specimen 1 exhibits slightly lower specific wear rate, therefore it is the more resistant to wear. This result is in accordance with its greater hardness when compared to the others as-deposited coatings. Specimens 2 and 3 displays similar wear volumes, which is


again attributed to their similar hardness profiles. Figure 5(b) reveals that specimens exposed during five and ten days at 500°C displayed similar specific wear rate to the non-heat-treated specimen. However, the specimen exposed during 20 days exhibits lower specific wear rate, therefore, it is the most resistant to wear. This last result is again in accordance with the higher hardness of this sample after the heat treatment. Table 4 Results of the linearisation analysis of samples in as-deposited and heat treated

conditions

Specific wear rate, k(mm3/N m)

Average STD Confidence interval 90% r2

Specimen 1 2.83 × 10–3 2.65 × 10–4 2.20 × 10–3 to 3.45 × 10–3 0.974 Specimen 2 3.44 × 10–3 3.78 × 10–4 2.55 × 10–3 to 4.33 × 10–3 0.965 Specimen 3 3.45 × 10–3 2.14 × 10–4 2.94 × 10–3 to 3.95 × 10–3 0.989 After 5 days at 500°C 3.69 × 10–3 5.57 × 10–5 3.56 × 10–3 to 3.83 × 10–3 0.999 After 10 days at 500°C 3.73 × 10–3 1.91 × 10–4 3.28 × 10–3 to 4.18 × 10–3 0.992 After 20 days at 500°C 2.58 × 10–3 3.18 × 10–4 1.83 × 10–3 to 3.33 × 10–3 0.956

Figure 6 Confidence interval for a confidence level of 90%, for the specific wear rate of specimens in as-deposited and heat treated conditions (see online version for colours)

In order to identify the wear mechanisms and to compare the worn surfaces with those observed in the surface of moulds, SEM examinations were carried out. Figure 7(a) displays a SEM micrograph of the surface of a coating of a mould after long time in service. From this picture is possible to see that the surface failure regions display pits and cracks. Considering the in-service conditions the surface damage probably was induced by synergetic effects of several mechanisms, namely thermal fatigue, corrosion and abrasion. However, pitting is the main active failure mechanism of the moulds. Figures 7(b) and 7(c) illustrates spherical cup depressions induced by the micro-scale abrasion tests on the specimen 3, before and after holding time for 20 days, produced after 500 ball rotations. Both scars display similar wear behaviour with two modes of wear: three-body abrasion (pitting) and two-body abrasion (grooving). The two-body abrasion mode is preferentially localised in the area of higher pressure (input slurry area)


and three-body abrasion mode is localised in the area of less pressure. Figure 8 shows these two modes of wear at higher magnification. The same wear mechanisms were observed in the other specimens in as-deposit and heat treated conditions. Although the scars displayed two modes of wear it can be stated that the region corresponding to three-body mode reproduces closely the wear mechanism present in the moulds surface. However, one should keep in mind that these results do not reproduce with accuracy the in-service condition of the moulds, since in service they are exposed to other damage mechanisms such as thermal and mechanical fatigue, oxidation and corrosion at high temperature.

Figure 7 SEM micrograph (a) of the surface of a coating of a mould after long exposure to high temperature, (b) scar with 500 rotations induced by the micro-scale abrasion tests on specimen 3, and (c) scar with 500 rotations induced by the micro-scale abrasion tests on specimen 3 after thermal exposure for 20 days (see online version for colours)

Figure 8 SEM micrograph of (a) three-body abrasion and (b) two-body abrasion


4 Conclusions

Micro-scale abrasion testing was used to study the interaction conditions between melted glass and the surface of the moulds. The influence of the arc current used in PTA process, as well as the effect of heat treatments, on the microstructure, hardness and wear resistance of the coatings was analysed. The increase in arc current increased the dilution of the base material changing the composition and microstructure of the deposits and reducing their hardness. Furthermore, the wear loss of material increased with increasing arc current. The micro-scale abrasion tests displayed two modes of wear: three-body abrasion (pitting) and two-body abrasion (grooving). Although the scars displayed two modes of wear it can be stated that the wear mechanism of the abrasion test reproduces with success the wear mechanism present in the moulds surface. The heat treatment performed on the sample 3 promotes increasing hardness of the coating with thermal exposure time. The enhanced hardness of the coatings is partially attributed to Ni-P precipitates on the grain boundaries. Specimen 3 heat treated for 20 days showed the best resistance to wear.

Future experiments will be devoted to study the high temperature abrasion resistance of the coatings, in order to reproduce the in service conditions of moulds.

Acknowledgements

The authors would like to thank the company ‘Intermolde’, for supplying the coated samples, and the Portuguese Foundation for the Science and Technology (FCT), through COMPETE programme from QREN and to FEDER, for financial support in the aim of the project number ‘013545’, as well as for the grant SFRH/BD/68740/2010.

References Balasubramanian, V., Lakshminarayanan, A.K., Varahamoorthy, R. and Babu, S. (2009)

‘Application of response surface methodolody to prediction of dilution in plasma transferred arc hardfacing of stainless steel on carbon steel’, International Journal of Iron and Steel Research, Vol. 16, No. 1, pp.44–53.

Chang, J.H., Chang, C.P., Chou, J.M., Hsieh, R.I. and Lee, J.L. (2010) ‘Microstructure and bonding behavior on the interface of an induction-melted Ni-based alloy coating and AISI 4140 steel substrate’, Surface and Coatings Technology, Vol. 204, No. 2, pp.3173–3181.

Cingi, M., Arisoy, F., Basman, G. and Sesen, K. (2002) ‘The effects of metallurgical structures of different alloyed glass mold cast irons on the mold performance’, Materials Letters, Vol. 55, No. 6, pp.360–363.

Ezugwu, E.O., Wang, Z.M. and Machado, A.R. (1998) ‘The machinability of nickel-based alloys: a review’, Journal of Materials Processing Technology, Vol. 86, Nos. 1–3, pp.1–16.

Fernandes, F., Lopes, B., Cavaleiro, A., Ramalho, A. and Loureiro, A. (2011) ‘Effect of arc current on microstructure and wear characteristics of a Ni-based coating deposited by PTA on gray cast iron’, Surface and Coatings Technology, Vol. 205, No. 16, pp.4094–4106.

Gant, A.J. and Gee, M.G. (2011) ‘A review of micro-scale abrasion testing’, Journal of Physics D: Applied Physics, Vol. 44, No. 7, 073001.

Gatto, A., Bassoli, E. and Fornari, M. (2004) ‘Plasma transferred arc deposition of powdered high performances alloys: process parameters optimisation as a function of alloy and geometrical configuration’, Surface and Coatings Technology, Vol. 187, Nos. 2–3, pp.265–271.


Guo, C., Zhou, J., Chen, J., Zhao, J., Yu, Y. and Zhou, H. (2011) ‘High temperature wear resistance of laser cladding NiCrBSi and NiCrBSi/WC-Ni composite coatings’, Wear, Vol. 270, Nos. 7–8, pp.492–498.

Gurumoorthy, K., Kamaraj, M., Rao, K.P., Rao, A.S. and Venugopal, S. (2007) ‘Microstructural aspects of plasma transferred arc surfaced Ni-based hardfacing alloy’, Materials Science and Engineering: A, Vol. 456, Nos. 1–2, pp.11–19.

Harsha, S., Dwivedi, D. and Agarwal, A. (2008) ‘Influence of CrC addition in Ni-Cr-Si-B flame sprayed coatings on microstructure, microhardness and wear behaviour’, The International Journal of Advanced Manufacturing Technology, Vol. 38, No. 1, pp.93–101.

Kesavan, D. and Kamaraj, M. (2010) ‘The microstructure and high temperature wear performance of a nickel base hardfaced coating’, Surface and Coatings Technology, Vol. 204, No. 24, pp.4034–4043.

Navas, C., Colaço, R., Damborenea, J. and Vilar, R. (2006) ‘Abrasive wear behaviour of laser clad and flame sprayed-melted NiCrBSi coatings’, Surface and Coatings Technology, Vol. 200, No. 24, pp.6854–6862.

Pouranvari, M. (2010) ‘On the weldability of grey cast iron using nickel based filler metal’, Materials & Design, Vol. 31, No. 7, pp.3253–3258.

Ramalho, A. (2010) ‘A reliability model for friction and wear experimental data’, Wear, Vol. 269, Nos. 3–4, pp.213–223.

Ramalho, A., Rodríguez, J. and Rico, A. (2011) ‘Microabrasion resistance of nanostructured plasma-sprayed coatings’, Int. J. Surface Science and Engineering, Vol. 5, No. 4, pp.250–260.

Sahoo, P. and Das, S.K. (2011) ‘Tribology of electroless nickel coatings – a review’, Materials & Design, Vol. 32, No. 4, pp.1760–1775.

Siva, K., Murugan, N. and Logesh, R. (2009) ‘Optimization of weld bead geometry in plasma transferred arc hardfaced austenitic stainless steel plates using genetic algorithm’, The International Journal of Advanced Manufacturing Technology, Vol. 41, No. 1, pp.24–30.

Annex D


Annex D

F. Fernandes, A. Ramalho, A. Loureiro, A. Cavaleiro, Mapping the micro-

abrasion resistance of a Ni-based coating deposited by PTA on gray cast iron, Wear,

292-293 (2012) 151-158.

Wear 292–293 (2012) 151–158


Wear

0043-16

http://d

n Corr

E-m

journal homepage: www.elsevier.com/locate/wear

Mapping the micro-abrasion resistance of a Ni-based coating depositedby PTA on gray cast iron

F. Fernandes n, A. Ramalho, A. Loureiro, A. Cavaleiro

CEMUC, Department of Mechanical Engineering, University of Coimbra, Rua Luıs Reis Santos, 3030-788 Coimbra, Portugal


Article history:

Received 2 February 2012

Received in revised form

28 May 2012

Accepted 30 May 2012Available online 7 June 2012

Keywords:

Mapping

Micro-scale abrasion

Three-body abrasion

Two-body abrasion

Hardfacing

Wear testing

48/$ - see front matter & 2012 Elsevier B.V. A

x.doi.org/10.1016/j.wear.2012.05.018

esponding author. Tel.: þ351 239 790 700; fa

ail address: [email protected] (F. Fer

a b s t r a c t

Micro-scale abrasion was used to characterize abrasion resistance of a nickel-based hardfacing alloy

deposited by Plasma Transferred Arc on gray cast iron using silica as abrasive agent. In order to

investigate the occurrence of different abrasion mechanisms, several test conditions were used,

namely: different silica abrasive contents and a range of normal loads and test durations. A scanning

electron microscope (SEM) was used to study the morphologies of the spherical cup-shaped depres-

sions induced by different test conditions. The results are discussed in terms of the effect of the

dominant wear mechanisms on the abrasion resistance and the influence of the test conditions on the

mechanism transition. The wear results lead to the conclusion that the specific wear rate essentially

depends on the wear mechanisms (rolling or grooving) involved and not on the tests conditions

employed, since these do not produce changes in the wear mechanism.

& 2012 Elsevier B.V. All rights reserved.

1. Introduction

The micro-scale abrasion wear test is a suitable experimentalmethod for evaluating and comparing the abrasive wear perfor-mance of a wide variety of materials. This technique has beensuccessfully applied to characterize metallic and non-metallic,bulk and coating materials, under the influence of a wide diversityof abrasive slurries [1,2]. It is well known that in the industry, themain wear problems in components are related to abrasion,which could co-exist, whether associated or not, with other formsof wear, such as adhesion or tribo-corrosion [3–5]. The usefulnessof the laboratory characterization depends on the ability of thetest conditions to reproduce real working conditions as closely aspossible.

In the micro-scale abrasion test, three distinct wear modeshave been identified, three-body or rolling abrasion, two-body orgrooving abrasion and the ridging wear mechanism [1,2,5–8].Additionally transitions between these wear modes could beobserved (rolling to grooving, grooving to ridging and ridging torolling); these are referred to as ‘‘mixed-mode’’ wear mechan-isms. Rolling abrasion is characterized by a hertzian type contactwhich occurs when the abrasive particles are able to freely rotatebetween the sphere and the material in test, producing multipleindentations, while grooving occurs if the particles remain fixedto the sphere and scratch the specimen surface. On the other

ll rights reserved.

x: þ351 239 790 701.

nandes).

hand, ridging wear mechanism is produced when the abrasiveparticles disappear in the contact region, the contact being madedirectly between the sphere and the specimen tested. Therefore,the ridging wear mode should be considered as transition fromabrasion to metal–metal sliding. Furthermore, according to someauthors [8,9], associated with the complex regimes above men-tioned, tribo-chemical interactions must be taken into account inthe wear calculations specially at high applied loads. This phe-nomenon is known to enhance wear resistance with an increasein applied load due to oxide formation on the surface.

In order to use micro-abrasion tests to evaluate the abrasivewear resistance of a system, either three-body or two-bodyabrasions need to be considered in the study; each wear modeis characterized by significantly different values of specific wearrate, which can differ by more than one order of magnitude [1,10].Some authors stated that, in micro-scale abrasion tests, the rollingabrasion wear mode leads to results that are more easily repro-duced [11,12]. Williams and Hyncica [13,14] discussed the pos-sibility of applying the continuum mechanics abrasion model toexplain the wear transitions. They verified that the quantitativemodels can be only applied to processes or materials where theboundary conditions are well established, which is not the casefor the micro-scale abrasion test. Nevertheless, Williams andHyncica suggested that the transition from rolling to groovingmust be a function of the critical thickness of the abrasive slurryin the contact. Following this suggestion, Adachi and Hutchings[6] used an indentation model to analyze the grooving wearconditions and they proposed a dimensionless parameter forthe severity of contact (S0), to be correlated with the transition

www.elsevier.com/locate/wear


dx.doi.org/10.1016/j.wear.2012.05.018

dx.doi.org/10.1016/j.wear.2012.05.018

dx.doi.org/10.1016/j.wear.2012.05.018


dx.doi.org/10.1016/j.wear.2012.05.018

Fig. 1. Typical microstructure of the as-deposited coating.

Fig. 2. Schematic diagram of the of the micro-abrasive wear equipment used.

F. Fernandes et al. / Wear 292–293 (2012) 151–158152

conditions. The ratio of hardnesses between the specimen and theball, Hs/Hb, was identified as the key parameter in the transitionfrom three-body to two-body abrasion. The results showed thatwith increasing severity of contact (S0) a transition from uniformthree-body to two-body abrasion occurs through a mixed threeand two body abrasion regime.

The technical literature reports several studies undertakenusing micro-scale abrasion equipment and describes how thedissimilar test conditions (volume fraction of abrasive in theslurry, abrasive material, applied load, ball and specimen materi-als, sliding distance and surface condition of the ball) influencethe dominant wear modes of different materials [1,6,10,15–18].The main objective of these studies is to track the dominant wearmodes created with different conditions to produce an associatewear-mode map, to serve as data to predict the conditions toensure either three-body or two-body abrasion. With regard tothe type of abrasive material to be used in micro-scale abrasiontests, the BS-EN 1071-6: 2007 standard recommends the use ofSiC F1200 abrasive; however, dissimilar abrasive materials withdifferent grain sizes have been used and, normally, the studiesalso report the use of diamond and Al2O3. It has been shown thatthe relative wear rates of the materials depended strongly uponthe abrasive type selected (abrasive shape and hardness) [19].Nevertheless, there are few references regarding the use of silicaas abrasive material. Because silica is the most common abrasivematerial, therefore responsible for the majority of practical abra-sion problems, it is important to study the test conditions whichproduce different wear mechanisms using silica abrasive slurry.

Thus, the purpose of this investigation was to study the testconditions which produce rolling and grooving wear modes withsilica slurry. An associate wear mode-map should be produced touse as data to ensure either one or the other abrasion wearmechanism’s intended further aim was to show that the specificwear rate essentially depends on the wear mechanisms (rolling orgrooving) involved and not on the test conditions employed, sincethese do not produce changes in the wear mechanism. The weartests were conducted in fixed ball equipment on a nickel-basedhardfacing alloy coating deposited by PTA – Plasma TransferredArc on gray cast iron. Regimes of wear behavior were identified byscanning electron microscope (SEM) and the results were ana-lyzed with a focus on the effects of wear mechanisms.


In this study, micro-scale abrasion equipment was used tostudy the influence of the test conditions on the dominant wearmodes of a nickel-based hardfacing coating. The coating analyzedwas produced by plasma transferred arc process (PTA) on a flatsurface of gray cast iron using a 128 A arc current. Fig. 1 showsthe typical microstructure of the as-deposited coating. As can beobserved, the coating has a relatively dense microstructure, freeof microcracks and few solidification voids with dendrites ofNi–Fe solid solution phase aligned along the direction of heatflow. Furthermore, it displays C-flakes (dark-floret like structures)randomly distributed in the matrix. The chemical composition ofthe materials, the conditions and the procedure adopted for thedeposition and the characterization of the coatings as well as theirmain physical structural and mechanical properties can be foundin a previous publication of the authors [20].

Fig. 2 represents a schematic diagram of the micro-abrasivewear equipment (fixed ball equipment). In this device, a specificnormal load is applied to press a rotating ball (located betweentwo-coaxial shafts) into the surface of the coated specimen(placed in a pivoted specimen holder) in the presence of a specificabrasive slurry. The abrasive slurry was continuously agitated by

a magnetic stirrer and dropped steadily onto the rotating ball toproduce a spherical cup depression. In this investigation, highpurity silica from Lusosilica, Lda (Portugal) company, with anaverage particle size of 3.1 mm and hardness around 1000 HV wasused as abrasive to produce spherical cup depressions in the flatsurface of the coated sample. As shown in Fig. 3(a), the silicaparticles are highly irregular and angular. The particle sizedistribution is plotted in Fig. 3(b). The slurry was prepared withdifferent silica contents (25, 50 and 65 vol%) in distilled water. AnAISI 52100 steel ball bearing, 25.4 mm in diameter and with ahardness of approximately 740 HV, was used as the counterpartto produce wear scars in the sample surface. The specimen testedhas a hardness of approximately 225 HV. The rotation speed ofthe ball was kept at 75 rpm (0.1 m/s of tangential speed) in all thetests. The coating was tested using three different normal loadvalues (0.1, 0.2 and 0.5 N) and five test durations (50, 80, 150, 200and 300 rpm), respectively 4, 6.4, 12, 16 and 24 m of slidingdistance. The test conditions are summarized in Table 1.

In order to produce an evenly prepared surface, the specimenwas polished with SiC abrasive paper up to 1000 grit before weartests. A new ball was used in the abrasion tests. The ball’s surfacewas prepared using the run-in procedure described by Gee et al.[21], and the orientation of the ball was changed randomly beforeeach test. Following the tests, the wear scars were examined by ascanning electron microscope (SEM), in order to discriminate thedissimilar wear modes. Abrasion maps of the material volumeloss as function of the sliding distance were plotted for each valueof normal load used. The results were discussed in terms of eitherthe tribological properties or the dominant wear mechanismsinvolved. The material wear volume was calculated by using Eq. (1).In this equation b represents the crater chordal diameter of the

Fig. 3. Silica abrasive particles. (a) Morphology and (b) particle size distribution.

Fig. 4. Results of the micro-scale abrasion wear test for the different test

conditions used.

Table 1Test conditions.

Ball material Steel AISI 52100

Ball diameter (mm) 25.4

Abrasive Silica SS40 (mean size: 3.1 mm)

Abrasive concentration 25, 50 and 65 vol% in H2O

Test duration (rotations) 50, 80, 150, 200, 300

Normal load (N) 0.1, 0.2, 0.5

Speed (rpm) 75

F. Fernandes et al. / Wear 292–293 (2012) 151–158 153

spherical cup depression produced and R is the ball radius used. Theb value was calculated using the average of two perpendicularmeasurements performed at the wear scar chordal diameter, withthe help of a Mitutoyo Toolmaker Microscope with x–y micrometertable. All the data was then collected in a single graph, andrepresented as function of the volume loss, with the productbetween the sliding distance and the normal load, to identify theconditions that produce similar wear mechanisms. To quantify thespecific wear rate of the material a linear approach to Archard’s lawwas used. This law is given by Eq. (2). In this equation V representsthe wear volume, k is the specific wear rate, N is the normal appliedload and x is the sliding distance. Moreover, a statistical analysis wasapplied to estimate the error of measurements, as described else-where [22].

V ¼ ðf� b4Þ=ð64� RÞ ð1Þ

V ¼ k� N � x ð2Þ

Eq. (1) allows the approximate calculation of the wear volumeof a spherical cup depression, with an implied assumption thatthe wear scar has a spherical geometry conforming to that of theball bearing used. In order to elucidate users, an analysis of theerror of Eq. (1) is given in Appendix A. Moreover, a correctionfactor is proposed to enhance the accuracy of this equation. Theanalysis carried out in Appendix A, showed that for b/R values upto 0.35 the equation assures an error lower than 1%.

3. Results

The micro-scale abrasion results for the different test condi-tions are shown in Fig. 4. Each graph plots the volume loss (V) asfunction of the sliding distance (x), for each normal load. Theresults of the tests performed with the lowest load (see Fig. 4(a))reveal that there are no effective changes in the wear volume withthe change in the concentration of abrasive slurry, especially forthe scar produced with low sliding distance. Inversely, for thehighest load (see Fig. 4(c)), the results show that the materialvolume loss increases with increasing concentration of abrasiveand the difference is sharper for higher sliding distances. The testsconducted with an intermediary load (see Fig. 4(b)), displayedmixed behavior. The scars produced with the lowest concentra-tion of abrasive slurry showed lower wear volume than thoseproduced with higher concentrations of abrasive slurry (50% and65% in volume), which displayed similar values of materialvolume loss. Moreover, it is observed that the wear volume


stabilizes with increasing sliding distance, whatever the concen-tration of abrasive used. This behavior could be associated withtwo effects: the running-in effect [23] and the variation of thecontact area/wear volume ratio throughout the duration of thetest. According to Blau [23], the running-in behavior is the netresult of simultaneous transitional processes occurring within theinterface, which may lead to non-linear evolution of the wear forshort sliding distances. Furthermore, at the start of the wear tests(with very small craters, consequently with small amounts of theradius of the spherical depression) it is well known that the areaof the crater grows much faster than the volume; at that point,the rate of pressure reduction is greater than the rate of change inthe volume, which may also justify the reduction in wear rate. Infact, the severity of contact index, proposed by Adachi andHutchings [6], is inversely proportional to the crater area; there-fore, a sudden increase in the area results in a reduction of theseverity index, which will induce a change in the wear rate. Thechanges in the wear volume described above are only partiallycorrelated with the test conditions, encouraging the investigationof the interactions between the abrasive slurry and the surface ofthe specimen in the contact zone. Scanning electron microscopy

Fig. 5. SEM morphologies of the worm surfaces of the sample performed with the com

and 4 m, (b) 0.5 N, 50%, 16 m and (c) 0.1 N, 25%, 16 m. Magnification of the wear mec

analyses were conducted on different spherical cup depressions inorder to identify the wear mechanism produced by the dissimilartest conditions.

Fig. 5 shows SEM micrographs of the dissimilar morphologiesof the spherical cup depressions induced by different test condi-tions of the micro-scale abrasion tests. As these images show, twodifferent wear mechanisms were identified, rolling abrasion (seeFig. 5(a)) and grooving abrasion (see Fig. 5(b)). Moreover, acombination of both wear mechanisms (see Fig. 5(c)) was alsoidentified for several contact conditions. In order to facilitateidentification of the wear mechanisms present in each wornsurface, magnified images are shown in Fig. 5(a1)–(c1). Rollingabrasion is characterized by multiple indentations in the wornsurface, (see Fig. 5(a1), whereas grooving abrasion (see Fig. 5(b1)is characterized by parallel grooves in the worn surface whichresults from the plastic deformation of the material. In turn, thesurface morphology, which displays a mixed-mode (rolling andgrooving abrasion), corresponds to a transition zone characterizedby a mixture of grooves and indentations, as shown in Fig. 5(c1).The transition of rolling to grooving abrasion occurs due to theembedding of particles in the surface. In this case grooving starts

bination of load, volume fraction of abrasive and sliding distance of: (a) 0.5 N, 50%

hanism present in the wear scar of: (a1) Figure a), (b1) Figure b), (c1) Figure c).

Fig. 6. (a) Scanning electron micrograph of the wear scar produced with 25% of

volume fraction of abrasive, load of 0.5 N and a sliding distance of 6.4 m: 1–ridging

wear mechanism zone, 2–grooving wear mechanism zone and 3–rolling wear

mechanism zone. (b) Magnification of the ridging wear mechanism zone.

Fig. 7. Wear modes observed in the wear scars produced by the different test

conditions.

Table 2Micro abrasion mechanism map as function of the

% of volume fraction of abrasives and loads used.


to occur in the zone of higher contact pressure, while rolling ismore commonly located in the surrounding area of less pressure,as shown in Fig. 5(c). Furthermore, in the tests performed with25 vol% abrasive and a load of 0.5 N, especially for the shortertests, the wear scars displayed areas with ridge wear mechanism,as Fig. 6 shows. Normally, this wear mechanism occurs for highvalues of normal load when the abrasive particles are absent fromthe contact region and contact is made directly between thesphere and the specimen tested, leading to ridging marks in thesurface [8], see zone 1 marked in Fig. 6(a). In this investigation theridging points will be disregarded since, in this type of wear, theabrasive particles do not influence the formation of the wear scarsand, therefore, the abrasion mechanism does not consider theeffect of the abrasive slurry.

% of volume fraction

of abrasiveLoad (N)

0.1 0.2 0.5

25 O // //

50 O y //

65 O y y

O–Rolling wear mechanism, y–Mixed-mode wear

mechanism and //–Grooving wear mechanism

4. Discussion

Considering the cumulative effect of the normal load multi-plied by the sliding distance (N� x) product, the wear map for thedifferent test conditions was plotted on a wear volume againstN� x graph by identifying the regimes of interactions based on

SEM observations. As can be seen in Fig. 7, three well definedzones (rolling abrasion zone, grooving abrasion zone and a mixed-mode wear) can be perceived. The graph reveals that differentcombinations of test conditions can produce the same wearmechanism. Moreover, it is observed that for some combinationsof concentration of abrasive slurry and load, the same wearmechanism is maintained even if the sliding distance is changed(conditions of (50 vol%; 0.1 N), (65 vol%; 0.1 N) and (25 vol%;0.2 N)). In these cases the wear volume always increases withincreasing sliding distance, even if the dominant wear mechanismis grooving or rolling abrasion. However, some conditions ofabrasive slurry concentration and load are more prone to induceone type of wear mechanism even when changing the slidingdistance. These changes are related to transitions in rolling tomixed-mode mechanism and mixed-mode to rolling wearmechanisms. For example, the wear scars produced with25 vol% abrasive using a load of 0.1 N produces both rollingabrasion and the mixed-mode wear mechanism depending onthe sliding distance. Table 2 summarizes the main wear mechan-isms produced by the different test conditions. The rolling,grooving and rolling–grooving abrasion regimes are differentiatedby using different symbols. The Table shows that low loads andhigh volume fraction of abrasive slurry enable three-body abra-sion, while grooving abrasion becomes stable at high loads andlow volume fraction of abrasive slurry. These last results are ingood agreement with previous studies [6,18].

Archard’s model is a suitable way to estimate a parameter ableto quantify the wear behavior of materials (the specific wearrate). However experimental results only fit the model well if asingle wear mechanism occurs in all the tests. The resultsdisplayed in Fig. 7 show that different combination of testconditions can produce the same wear mechanism. So, in thisstudy the experimental results were organized into two groupsdepending on the wear mechanism involved; rolling or grooving.

Fig. 8. (a) Results of the micro-scale abrasion wear test for the rolling and

grooving wear points and (b) error bars of the specific wear rate for a confidence

interval of 90%.

Table 3Results of the linearization analysis of the rolling and grooving zones.

Specific wear rate, k (mm3/N m)

Average STD Confidence interval 90% r2

Rolling zone 5.56�10�3 7.15�10�4 4.03–7.09�10�3 0.812

Grooving zone 7.79�10�4 8.21�10�5 6.20–9.39�10�4 0.937


Subsequently, Archard’s model was applied to quantify thespecific wear volume characteristic of each wear mechanism.

The micro-scale abrasion test results of the rolling and groov-ing abrasion points are shown in Fig. 8, as is the material volumeloss function of the product of the sliding distance by the normalload. For each set of data, a straight linear trend was fitted to thepoints, obtaining the specific wear rate as the slope of therespective linear trend line, as shown by Eq. (2). A completereliability analysis of the specific wear rate is summarized inTable 3 and, for each case, the confidence bands are plotted inFig. 8(b) for a confidence level of 90%. The results show that thespecific wear rate obtained from the rolling experimental datapoints is around one order of magnitude higher than in thecorresponding grooving points. In fact, the rolling wear mode isa more effective wear mechanism than the grooving mode,leading to higher values of specific wear rate [5,18,24]. Further-more, the reliability of results indicates that the correlationcoefficient for the specific wear rate calculated from the rollingand grooving points is high, meaning that the quality of the linearfittings displayed in Fig. 8 is suitable. Therefore, according to

these results, the specific wear rate can be estimated using pointsproduced with different test conditions, since a single wear modeis sustained. Furthermore, these results lead to the conclusionthat the specific wear rate essentially depends on the wearmechanisms (rolling or grooving) involved and not on the testsconditions employed, since these do not produce changes in thewear mechanism.

5. Conclusion

(i)
In this investigation, silica slurry was used as an abrasive inmicro-scale abrasion equipment to study the influence ofdifferent test conditions on the dominant wear modes of aNi-based coating deposited by PTA on gray cast iron.
(ii)
The effect of different concentrations of abrasive slurry,sliding distances and normal loads on the micro-abrasion ofa Ni-based coating were investigated and showed significantdifferences between micro-abrasion rates and wear mechan-isms as a function of these parameters.
(iii)
The micro-abrasion mechanism maps were represented interms of volume loss as a function of the sliding distancemultiplied by the normal load. The results showed thatdifferent combinations of tests can produce the same wearmechanism. Low values of normal load and high contents ofabrasive particles enable rolling abrasion, while groovingabrasion is stabilized at high loads and low volume fractionof abrasive slurry.
(iv)
The micro-scale abrasion results involving rolling and groov-ing abrasion data, based on calculations of the specific wearrate, revealed that this parameter can be successfully calcu-lated using data points produced by dissimilar test condi-tions since they give rise to a single wear mode. This resultled to the conclusion that the specific wear rate essentiallydepends on the wear mechanisms involved and not onthe test conditions employed, since a single wear mode issustained (grooving or rolling abrasion).
(v)
Abrasion by rolling wear induced a specific wear rate aroundone order of magnitude higher than the value obtained undergrooving abrasion.
Acknowledgments

The authors wish to express their sincere thanks to thecompany ‘‘Intermolde’’, for supplying the coated samples, andthe Portuguese Foundation for the Science and Technology (FCT),through COMPETE program from QREN and to FEDER, for financialsupport in the aim of the project number ‘‘013545’’, as well as forthe grant (SFRH/BD/68740/2010).

Appendix A. Error analysis of the equation of the wearvolume

This section aims to elucidate authors about the resultant error byusing an approximate equation to calculate the volume of a sphericalcup depression induced by the micro-scale abrasion tests.

The original equation used to calculate the volume of thespherical cup depression is given by the Eq. (A1), where b and h

represent respectively the crater chordal diameter and the depthof the spherical cup depression (see Fig A.1). However, in micro-scale abrasion testing an approximate equation of Eq. (A1) iscurrently used (see Eq. (A2)). In this equation R represents the ballradius of the ball bearing used. The error of using an approximate

Fig. A.2. Error associated to the approximate equation to estimate the volume of a

wear scar without and with corrective factor..

Fig. A.1. Schematic diagram of a crater induced by the micro-scale abrasion tests..


equation can be calculated by using Eq. (A3).

V ¼ ð1=6Þ � f� h� ð3� ðb=2Þ2þh2Þ ðA1Þ

Vn¼ ðf� b4

Þ=ð64� RÞ ðA2Þ

error¼ ðVn�VÞ=V ðA3Þ

It is evident from Fig A.1 that increasing the depth h of thespherical cup depression increases the value of the chordaldiameter b of the wear scar, until a maximum value of b equalto 2R is reached. So, discretizing the value of b in small intervals itis possible to calculate the associated values of V and Vn for eachrelation between b and h. The value of h (given by Eq. (A4)), can beeasily deduced by subtracting the value of c from the value of R, asrepresented in Fig A.1. The value of c is obtained by applyingPythagoras’ Theorem at the triangle. Applying Eq. (A3) at thedissimilar combinations of b and h it is possible obtain the errorassociated by using an approximate equation to estimate thevolume of a spherical cup depression induced by the micro-scaleabrasion test. The results of the error are plotted in Fig A.2, dashedline, as function of the ratio between the crater’s chordal diameterand the radius of the ball with the relative error associated. As canbe seen in this figure, the resulting error is very small for lowratios of b/R, however with an increase in this ratio, the errorincreases. For example considering a spherical cup depressionwith a chordal diameter equal to the radius of the ball, the errorcommitted in the calculations is about 10 percent.

h¼ R�ðR2�ðb=2Þ2Þ1=2

ðA4Þ

The level of correlation between the ratios of the depth of thespherical cup depression and the ball radius with the error proved tobe linear. So, according to this relationship it is possible to establish a

correction factor for a desired value of error. For example, for adesired error less than 1% a correction factor can be calculated asindicated by Eq. (A5). By manipulating the equation it is possible todetermine the limits as a function of the ratio between b and R (seeEq. (A6)). Fig A.2 displays the error obtained when the approximateequation to estimate the volume of the wear scar is used, correctionfactor applied (continuous lines) and without the correction factor(dashed lines). As can be seen in the figure, the error of using theexpression with the corrective factor is one hundred times smallerthan that obtained with the expression without the corrective factor.Although, the correction factor gives more accurate results for lowratios between the chordal crater and the radius of the ball bearing,the approximate equation to evaluate the volume of a spherical cupdepression is appropriate, due to the small error induced.

CF¼1 ( ðh=RÞr0:015

1=ð1�0:66� ðh=RÞÞ ( 0:015r ðh=RÞr0:55

(ðA5Þ

CF¼1 ( ðb=RÞr0:345

1=ð1�0:66� ð1�ð1�ðb=ð2� RÞÞ2Þð1=2ÞÞÞ ( 0:345rðb=RÞr1:786

(

ðA6Þ

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[2] A.J. Gant, M.G. Gee, A review of micro-scale abrasion testing, Journal ofPhysics D: Applied Physics 44 (2011) 073001.

[3] M.F.C. Andrade, R.P. Martinho, F.J.G. Silva, R.J.D. Alexandre, A.P.M. Baptista,Influence of the abrasive particles size in the micro-abrasion wear tests ofTiAlSiN thin coatings, Wear 267 (2009) 12–18.

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[6] K. Adachi, I.M. Hutchings, Wear-mode mapping for the micro-scale abrasiontest, Wear 255 (2003) 23–29.

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[17] M.M. Stack, M.T. Mathew, Mapping the micro-abrasion resistance of WC/Cobased coatings in aqueous conditions, Surface and Coatings Technology 183(2004) 337–346.

[18] A. Ramalho, J.R. Perez, A.R. Garcia, Microabrasion resistance of nanostruc-tured plasma-sprayed coatings, International Journal of Surface Science andEngineering 5 (2011) 250–260.

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[20] F. Fernandes, B. Lopes, A. Cavaleiro, A. Ramalho, A. Loureiro, Effect of arccurrent on microstructure and wear characteristics of a Ni-based coatingdeposited by PTA on gray cast iron, Surface and Coatings Technology 205(2011) 4094–4106.


[21] M.G. Gee, A.J. Gant, I.M. Hutchings, Y. Kusano, K. Schiffman, K.V. Acker,S. Poulat, Y. Gachon, J.v. Stebut, P. Hatto, G. Plint, Results from an inter-laboratory exercise to validate the micro-scale abrasion test, Wear 259(2005) 27–35.

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Annex E


Annex E

F. Fernandes, T. Polcar, A. Loureiro, A. Cavaleiro, Room and high temperature

tribological behavior of Ni-based coatings deposited by PTA on gray cast iron,

(2014), under review, “Tribology International”.

Annex E


Room and high temperature tribological behavior of Ni-based coatings deposited by

PTA on gray cast iron

F. Fernandes1,*

,T. Polcar2,3

, A. Loureiro1, A. Cavaleiro

1

1CEMUC - Department of Mechanical Engineering, University of Coimbra, Rua Luís Reis

Santos, 3030-788 Coimbra, Portugal.

2Department of Control Engineering Czech Technical University in Prague Technicka 2,

Prague 6, 166 27 Czech Republic.

3n-CATS University of Southampton Highfield Campus SO17 1BJ Southampton, UK.

*Email address: [email protected], tel. + (351) 239 790 745, fax. + (351) 239 790

701

Abstract

In the present investigation the effect of the substrate dilution on room and high

temperature (550 and 700 ºC) tribological behavior of nickel based hardfaced coating

deposited by plasma transferred arc onto a gray cast iron was investigated and compared to

the uncoated gray cast iron. At room temperature, the wear loss of coatings was independent

of the substrate dilution and similar to the gray cast iron. At high temperatures, coating

produced with high dilution displayed the highest wear resistance between all the samples.

This is attributed to the formation of a protective tribo-layer resulting from the agglomeration

of a high amount of oxide debris due to its lower oxidation resistance when compared to the

sample produced with low dilution.

Keywords: Plasma transferred arc, Ni-based alloy, Substrate dilution effect, Tribology, High

temperature wear

1. Introduction

Cast irons and copper-alloys are commonly used in the production of glass molds and

accessories for glass industry, owing to their excellent thermal conductivity and relatively low

cost [1]. Molds are often exposed to severe conditions of abrasion, oxidation, wear, fatigue at

high temperature, due to repeated contact with melted glass, causing deformation or failure of


Annex E


parts and, thus, compromising the product quality and increasing the maintenance costs. To

overcome this shortcoming protective coatings are normally applied at molds surfaces with

the aim of increasing their lifetime in harsh environments [2, 3]. A wide diversity of coating

materials have been successfully applied to protect the surface of these components, such as:

cobalt, iron and nickel alloys [2, 4, 5]. The latter have been especially used due to their

outstanding performance under extreme high temperature conditions [5-8]. Several hardfacing

and thermal spraying processes have been used to coat the molds: (i) plasma transferred arc

(PTA) and gas tungsten arc welding (GTAW), (ii) flame spraying (FS), high velocity oxy fuel

(HVOF) and atmospheric plasma spraying (APS). Hardfacing processes have been

preferentially used in the protection of mold surfaces because they promote a strong

metallurgical bonding between the coating and the substrate, condition required to achieve

high quality adherent coatings and to avoid the catastrophic failure in service, as opposed to

thermal spraying processes where only a mechanical bonding is established. Among the

hardfacing processes referred to above, PTA has been widely used to protect the surface of

molds, since it produces high quality thick coatings, offering both optimal protection with

minimal thermal distortion of parts and high deposition rates in single layer deposits [4, 5].

However, the properties and the quality of deposits are strongly dependent on the dilution of

the substrate promoted by the PTA process. Low dilution provides coatings with similar

chemical composition to the added metal powder, condition for achieving enhanced

mechanical properties, wear and oxidation resistance [9]. To avoid adhesion problems and,

therefore, not to compromise the performance of molds in service, it is common to increase

the dilution through the change of the most relevant deposition parameters even if the

mechanical properties, wear and oxidation resistance are diminished. In our previous studies

[10, 12] the effect of increasing dilution, promoted by change of arc current on the structure,

mechanical properties, oxidation resistance and wear behavior of a nickel-based alloy

deposited onto a gray cast iron have been investigated. Regarding the wear behavior of

coatings, ball cratering wear tester has been selected to reproduce the wear mechanism

produced by the interaction between the melted glass and the surface of molds (three body

abrasion). However, this equipment does not allow consider the high temperature effect.

Hence, as a way to complement the previous studies, the aim of this investigation is to study

the effect of increasing substrate dilution on the tribological performance at room and

essentially at high temperature of as deposited Ni-based coating deposited onto gray cast iron,

using a pin on disc tribometer apparatus. Comparison of these results with those achieved for

the base material is also provided.

Annex E



A nickel-based hardfacing alloy powder named Colmonoy 215 (from Colmonoy

company) was deposited by plasma transferred arc using a Commersald Group ROBO 90

machine onto flat surfaces of gray cast iron blocks, currently used in the production of molds

for the glass industry. Two different arc currents (100 and 140 amperes) were used on the

depositions to produce coatings with two dissimilar levels of dilution. Hereinafter, the

coatings deposited using 100 and 140 amperes will be designated as C100 and C140,

respectively. Table 1 gives the optimized parameters used in the depositions, while Table 2

depicts the nominal chemical composition of the base material and Colmonoy 215 powder.

Before coating, the substrate blocks were preheated at 480 ºC to reduce the cooling rate and

therefore to avoid the susceptibility to cracking during deposition. Using the same deposition

conditions, several weld beads, each with approximately 3 to 5 mm thick, were deposited in

parallel on the cast iron, ensuring some overlapping between them in order to obtain a 35×35

mm coated surface area and a quasi-flat surface. Specimens containing cast iron and coating

were removed from each block for microstructural and micro-hardness analysis. Further,

specimens with dimension of 20×20×5 mm were removed from each block for tribological

testing. The coated surfaces were first machined to provide flatness and then grinded and

polished following standard procedures to obtain a uniform roughness of about 0.51 – 0.62

µm.

The tribological behavior of the coatings and gray cast iron was evaluated at room and

high temperatures using a pin-on-disk tribometer. As the operating surface temperature of

molds is currently in the range of 500 - 900 ºC, temperatures of 550 and 700 ºC were selected

to perform the experiments. Although in the real working conditions of molds glass acts as an

abrasive material, due to its melted state at high temperature, for the tribological tests we

have selected as counterpart a harder and more resistant material: Al2O3 balls. Thus we ensure

that the wear volume loss of coatings will be overestimated in relation to the expected loss in

actual service conditions. All the wear tests were conducted at a constant linear speed of 0.10

m.s-1

, a load of 5 N, a relative humidity of 48±5%, and 5000 cycles. The radius of the wear

tracks was set to 5.3 mm. During the tests, the friction coefficient was continuously recorded.

Two tests were performed under identical conditions for each coating. After tribological tests,

the wear rate of the coatings was determined from the cross section wear track area using a

3D optical profilometer and the wear tracks and wear debris were characterized by scanning

Annex E


electron microscopy with energy dispersive X-ray spectroscopy (SEM-EDS) and by Raman

spectroscopy.

Table 1 – Main deposition parameters.

Main arc

current

(A)

Powder

feed rate

(rpm)

Travel

speed

(mm/s)

Powder

feed gas

flow rate

(l/min)

Plasma

gas flow

rate

(l/min)

Shielding

gas flow rate

(l/min)

Torch

work

distance

(mm)

Oscillation

(mm)

Preheat

temperature

ºC

C100 100 20 2 2 2.2 20 13 4 480

C140 140 20 2 2 2.2 20 13 4 480

Table 2 - Nominal chemical composition (wt.%) of the gray cast iron and nickel-based alloy powder.


Gray cast iron 3.60 0.60 2.00 < 0.20 < 0.04 < 0.20 <0.50 0.50 0.10 0.20 Balance

Hardfacing C Cr Si B Fe Al F Co Ni

Ni-alloy 0.14 2.45 2.56 0.86 1.08 1.30 0.01 0.08 Balance


3.1. Microstructure, composition and hardness of the as deposited coatings

The typical microstructure of gray cast iron and coatings produced with 100 and 140 A

is depicted in Fig. 1. Both coatings display relatively dense microstructure with excellent

adhesion to the substrate, free of microcracks and only a few solidification voids. Their

typical microstructure consists of dendrites of Ni-Fe solid solution phase aligned along the

direction of heat flow. Flakes of carbon (dark-floret like phase) randomly distributed in the

matrix can be perceived too, being more notorious on C140 coating. The high amount of dark-

floret phase on C140 coating is in good agreement with the observed increase of dilution with

increasing arc current. Following the procedure described in reference [11], dilutions of ~28

and ~59% were measured for C100 and C140 coatings, respectively. The dilution is beneficial

in terms of improving the adhesion, but it can bring some disadvantages, such as change in

the hardness and chemical composition, and lowering of wear and oxidation resistance of the

coatings [9]. The chemical composition of coatings, assessed by SEM-EDS given in Table 3,

were evaluated from the average of several measurements performed on an area 400×400 µm

in the cross section at middle thickness of each coating. This procedure was used to ensure

that the chemical composition was determined in a representative volume of material in that

zone, thus avoiding problems that punctual analyses can give in heterogeneous materials. The

Annex E


results show that the main differences were in the iron, silicon and carbon contents, in good

agreement with the different dilutions in the samples. The gray cast iron shows flakes of

graphite evenly distributed in a ferritic matrix.

Figure 1 - Optical micrograph of: a) C100 coating, b) C140 coating, c) interface between coating and substrate of

C140 coating, d) gray cast iron.

Table 3 - Nominal chemical composition (wt.%) of coatings evaluated by SEM-EDS.

wt. % of elements

Coating Ni Fe C Si Al Cr

C100 85.40 8.47 0.30 2.17 0.81 2.86

C140 55.71 37.73 1.26 2.52 0.62 2.16

Fig. 2 shows the Vickers hardness profiles across the interface of the deposited

coatings, measured with 5 N applied load. As it would be expected, the increase in the weld

current, and therefore increase in dilution, gives rise to a decrease of the hardness either in the

coating, undesirable in terms of wear, or in the heat affected zone (HAZ), which is favourable

in terms of toughness of this region. The average surface hardness of coatings C100, C140

and gray cast iron was 300, 195 and 150 Vickers (HV0.5), respectively. The high hardness of

the HAZ zone in relation to the non-thermal affected based material is related to the formation

of hard and brittle phases. The detailed characterization of the microstructure of each coating

a) b)

c) d)

Graphite

Grains of ferrite

Coating

Gray cast iron

Graphite

Graphite

Annex E


and the hardness variations with increasing dilution can be found in a previous publication of

the authors [11].

Figure 2 - Hardness profile across the interface of C100 and C140 coatings.

4. Tribological behavior of coatings

4.1. Wear rate

Figure 3 shows the wear loss rate of coatings and gray cast iron tested at different

temperatures. It clearly shows that at room temperature the wear loss of coatings and gray cast

iron is very low, and almost independent of the substrate dilution. C100 coating displays a

little higher wear resistance than C140 coating, in agreement with its higher hardness. The

observed similar level of wear rate of the gray cast iron as compared to the Ni coatings at RT

is due to its higher amount of graphite that can act as solid lubricant, leading to a less volume

loss of material. It should be pointed out that the low hardness displayed by the gray cast iron

compared to the coated samples is compensated by the presence of graphite on the sliding

contact [13, 14].

At temperature of 550 ºC an abrupt increase of the wear rate of coatings was observed;

however, it dropped with further increase to 700 ºC. The first increase of the wear rate can be

attributed to the spontaneous oxidation, while, as it will be seen later, following decrease of

wear loss at 700 ºC could be related to the formation of high amount of oxides on the wear

track preventing further wear of the coatings. Besides, coating produced with high dilution

displayed much higher wear resistance at elevated temperatures than C100 coating. This

-3 -2 -1 0 1 2 3 4 5 6100

150

200

250

300

350

400

450

500

C100 coating

C140 coating

PMZ

Base material

HAZCoating

Fu

sio

n L

ine

Hard

ness H

V0

.5


Annex E


behavior is a result of the low oxidation resistance of C140 coating that will contribute to the

formation of a high amount of oxides in the contact zone. A continuous increase of the wear

rate with increasing temperature was observed for the gray cast iron. This material showed

much higher wear rate than the Ni-based coatings, indicating that the latter should perform

better at high temperature. This is an expected result since rapid loss of the material

mechanical strength and increasing oxidation rate with temperature are typical of gray cast

iron.

Figure 3 - Variation of the wear rate of coatings and gray cast iron with test temperature.

4.2. Friction coefficient

The evolution of the friction coefficient of coatings and gray cast iron with increasing

test temperature is shown in Figure 4. As expected, all the friction curves displayed two

distinct regions: running-in stage and steady state stage. At RT, the running-in stage for both

Ni-based coatings is similar and characterized by strong increase of the friction coefficient to

the highest value of 0.74, followed by rapid decrease down to 0.41, after the first 1000 laps.

On the other hand, in gray cast iron, friction coefficient slightly increased over time reaching

the steady state after 4000 laps with a friction coefficient of about 0.39. Gray cast iron

exhibited relatively low friction coefficient at RT as compared to the Ni-based alloys.

According to the literature [13, 14], this behavior is related to the presence of graphite on the

worn surface that acts as a solid lubricant compensating the lower hardness displayed by the

material and thus leading to wear rate values similar to Ni-based coatings.

0

50

100

150

200

250

300

350

400

450 RT

550 0C

700 0C

Cast ironC 140C 100

Wear

rate

(10

-6 m

m3/N

.m)

Annex E


A long running-in period was also noticed with a progressive increase of the friction

coefficient for the gray cast iron when the temperature was 550 ºC; however, friction was

much higher suggesting different contact conditions. Ni-based coatings displayed very short

running-in periods with friction stabilized at similar values as in RT tests. Finally, the tests at

700 ºC exhibited different trends between cast iron and coatings. In the first case, a small

decrease in the final friction coefficient value was achieved whereas an inverse situation was

observed for coated samples, with friction stabilized at values of approximately 0.6 and 0.65

for C100 and C140 coatings, respectively. Such trends suggest that different contact

conditions should have been taking place in both situations.

In many cases, especially at high temperatures, the friction curves displayed quite

unstable values. According to Kesavan and Kamaj [15], such behavior demonstrates the

presence of tribolayer on the sliding surfaces; the tribolayer is continually worn out and

formed again, which results in short-term friction coefficient oscillations.

Figure 4 - Friction coefficient evolution with increasing test temperature of: a) C100 coating, b) C140 coating, c)

gray cast iron.

0 1000 2000 3000 4000 50000.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

RT

550 0C

700 0C

a)

C100 Coating

Friction c

oeffic

ient

Number of cycles

Annex E


Figure 4 (continued).

4.3. Wear mechanisms and surface analysis of coatings

The changes of the wear rates and friction coefficients described above can be

attributed to a great number of factors, such as the chemical composition, physical properties

of materials, or applied test conditions. Therefore, the investigation of the interaction between

the specimen-counterpart pair was carried out. Scanning electron microscopy with energy

dispersive X-ray spectroscopy (SEM-EDS) and Raman spectroscopy were used to

characterize the dominant wear mechanisms and the wear debris originated by the tribological

testing.

0 1000 2000 3000 4000 50000.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

RT

550 0C

700 0C

b)

C140 Coating

Friction c

oeffic

ient

Number of cycles

0 1000 2000 3000 4000 50000.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

RT

550 0C

700 0C

c)

Gray cast iron

Friction c

oeffic

ient

Number of cycles

Annex E


4.3.1. Worn surfaces at room temperature

Typical SEM micrographs in secondary electrons of the worn surface of coatings and

gray cast iron tested against Al2O3 balls at RT are depicted in Figure 5. The wear tracks of Ni-

based coatings are similar; they present a rough morphology with evidence of tribo-oxidation.

The oxide adhered layer is cracked, which reveals its brittle nature. Raman spectra performed

at the oxidized worn surface of the coatings (see spectra a) and b) in Fig. 6 for C100 and C140

coatings, respectively) confirmed the presence of the oxides, namely Raman active modes

assigned to Ni-O [16] and C [17]; the latter almost inexistent in C100 sample spectrum as a

consequence of its much lower dilution. These results are in good agreement to EDS analysis

(see Fig. 5 d) for C100 coating as an example), which showed essentially strong signals of Ni

and O; vestiges of Fe, Si, Al and C were also detected. Similar wear debris was found on the

wear scars of the balls too. During the wear tests the friction, generated by the sliding of the

ball against the coating surface, causes the increase of the local temperature, mainly in the

protuberances contact, promoting the tribo-oxidation of the surface. During this phase, the

metal-ceramic contact will progressively decrease with more and more surface being covered

by an oxide phase. With the test running, the oxide tribo-layer will grow and the contact shall

be governed between this growing tribo-layer and the counterpart, until a steady state will

occur. Metal-ceramic contact gave rise to a higher friction coefficient that should decreases

until a constant and lower value corresponding to the same oxide, NiO as shown by Raman

spectroscopy, formed on the top surface. It explains the similar level of wear resistance and

friction values of the coatings even despite of their discrepancy in hardness and chemical

composition due to different dilution levels. The high hardness of the metallic coatings, in

conjunction with the protective effect of the hard NiO oxide, led to very low wear volumes.

Distinct wear behavior was found for the gray cast iron. Its wear track displays a very

clean surface with fine scratches parallel to the relative sliding movement (suggesting

abrasion wear mechanism) and some adherent wear debris. Further, in some places, black

regions can be noticed, which can be attributed to either the presence of graphite or its pull-

out from the material [18]. EDS analysis performed at the wear debris revealed that they

represent a mixture of original and oxidized material, including a weak signal from graphite

lamellas. Similar phases were detected on clean worn surface, although C peak was much

more intense there. Raman spectroscopy analysis showed the presence of hematite (α-Fe2O3)

[19] and nanocrystalline graphite through the characteristic D and G bands [20] (see spectrum

c) and d) of Fig. 6 for wear debris and clean surface, respectively). During sliding

Annex E


experiments the movement of the ball promotes the wear of surface specimen with production

of both oxide and graphite debris. Graphite will act as a solid lubricant in tribo-film,

promoting a decrease of friction coefficient and having, indirectly, an important impact on the

wear [13, 14]. As a consequence, similar level of wear resistance was observed between base

material and coatings, in spite of the lower hardness of gray cast iron.

Figure 5 - SEM morphology in secondary electrons of worn surface of: a) C100 coating, b) 140 coating, c) cast

iron; all tested at room temperature. d) SEM/EDS spectra of point 1 marked in Fig. 5 a). e) and f) wear track

profiles of: C140 coating, and gray cast iron, respectively.

1

a)

b)

c)

Wear debrisBlack

regions

0

0.5

1

1.5

2

2.5

3

3.5

0.1 0.2 0.3 0.4 0.5

Depth

(µ

m)

Scan length (µm)

0

0.5

1

1.5

2

2.5

3

3.5

0.1 0.2 0.3 0.4 0.5

Depth

(µ

m)

Scan length (µm)

e)

f)

0.4 0.8 1.2 1.6KV

Inte

nsity (

a.u

.)C Kα

O Kα

Ni Lα

Fe

Lα

Al Kα

Si Kα

d)

Annex E


Figure 6 - Typical Raman spectra performed on the worn surface of: a) C100 coating tested at RT, b) C140

coating tested at RT, c-d) adherent wear debris and clean surface of gray cast iron, respectively, tested at room

temperature.

4.3.2. High Temperature

When sliding occurred at 550 ºC, C100 and C140 coatings also showed oxides on the

contact surface. However, C140 wear track was almost fully covered whereas C100 exhibited

an irregular surface with a mixture of clean wear zones and oxidized ones with signs of

adhesive and abrasive wear (see Fig. 7 a-b), suggesting an incomplete formation of an oxide

transfer layer. The difference between the two samples can be interpreted considering their

oxidation behavior. In fact, C100 sample oxidizes through the formation of a protective

mixture of oxides based on B / Si over a Ni-based oxide layer [10]. Although these oxides are

protective in static oxidation conditions, B-oxide melts below 600 ºC [21, 22], which leads to

its easy removal hindering the formation of a stable oxide layer over the surface. As the

process requires some time, since temperatures are still low being oxidation rates reduced

[10], the destruction of the growing oxide layers is relatively easy and the formation of a

transfer layer is limited due to the dragging effect caused by the ball sliding and melted B-

oxide. Therefore, high wear volumes are expected and indeed the steep increase in the wear

rate can be noticed in Fig. 3. In the case of C140 sample, besides Ni-based oxides, iron oxide

is also formed during oxidation at these temperatures [10]. The oxidation rates are much

higher compared to C100 sample due to higher dilution. Therefore, a higher amount of oxide

debris can be produced facilitating the formation of a continuous oxide transfer layer. The

surface is thus protected quickly and, in spite of the much higher wear rate in comparison to

200 400 600 800 1000 1200 1400 1600 1800 2000

C

b)

d)

c)

a)

Fe2O

3NiO

Re

lative inte

nsity

Raman shift [cm-1]

Annex E


RT test, the wear resistance is higher than that for C100 sample. Raman analysis performed at

the adhesive wear debris on C100 coating indicated that they were mainly Ni-O, although,

according to the EDS analysis, Fe-O, Si-O and Al-O should also coexist in much less amount.

It is in good agreement to our previous work on the oxidation behavior of this coating [10]. At

the clean zone (bottom part) of the worn surface EDS showed non-oxidized coating material.

For C140 coating, Raman analysis showed similar oxides but, in this case, a significant

amount of iron oxides (hematite and magnetite (Fe3O4)) were also detected, in good

agreement to higher level of dilution and to the oxidation studies where this oxide was found

to be the main constituent of the oxide scale [10]. Additionally, strong signals of C-based

phases were detected by EDS.

Gray cast iron was fully covered with an oxide layer at this temperature, with

scratches parallel to the movement of the sliding. The main oxides identified in the wear track

were hematite and magnetite (see Raman spectrum c) of Figure 8). The softening of the

material, induced by the temperature increase, together with the total depletion of graphite

from the track surface justify the sharp increase in the wear rate, as well as the increase in the

friction coefficient in comparison to results at RT.

Annex E


Figure 7 - SEM morphology in secondary and backscattered electrons of worn surface of: a-b) C100 coating, c-

d) 140 coating, e-f) cast iron; all tested at 550 ºC.

Wear

debris

a)

c)

e)

d)

f)

b)

http://en.wikipedia.org/wiki/Backscatter

Annex E


Figure 8 - Raman spectra performed on the worn surface of: a) C100 coating tested at 550 ºC, b) C140 coating

tested at 550 ºC, c) gray cast iron tested at 550 ºC.

At 700 ºC the oxidation rate of Ni-based coatings increased significantly allowing

rapid and easy formation of a continuous oxide layer on the worn surfaces (see Fig. 9). The

oxide acts as a protective layer decreasing the wear loss of coatings. As the major wear

volume loss is produced during the formation of the protective oxide layer, faster production

of tribolayer decreased the wear rate in relation to 550 ºC testing (see Fig. 3). Similar results

were achieved and in detail explained by Kesavan et al. [15]. Furthermore, since the oxidation

rate of C140 coating is higher than that of C100 coating, the latter will take more time to form

the oxide layer generating a higher coating wear volume in the initial stage of the wear

process. As a consequence, total wear rate after the test will be higher as well. Raman analysis

of the wear tracks of both coatings showed the presence of Ni-O (see spectra a) and b) of Fig.

10); moreover, iron oxides were identified in C140 coating.

On the other hand, gray cast iron oxidizes easily but the oxide does not support the

sliding process and does not protect efficiently the base material. The increase in the

temperature promotes a severe softening of the material and a high oxidation rate leading to a

strong abrasion and, therefore, a very high wear loss. Raman analysis confirms the extensive

formation of hematite in the wear track (spectra c of Fig. 10); signs of magnetite were also

detected.

In summary, although dilution reduces the hardness of coatings, it has beneficial effect

at high temperature tribological behavior due to the faster formation of an oxide layers (owing

to its less oxidation resistance) which protect the coating surface against wear.

200 400 600 800 1000 1200 1400 1600 1800 2000

Fe3O

4

b)

c)

a)

Fe2O

3NiO

Re

lative inte

nsity

Raman shift [cm-1]

Annex E


Figure 9 - SEM morphology in secondary and backscattered electrons of worn surface of: a-b) C100

coating, c-d) 140 coating, e-f) cast iron; all tested at 700 ºC.

a)

c)

e)

d)

f)

b)

http://en.wikipedia.org/wiki/Backscatter

Annex E


Figure 10 - Raman spectra performed on the worn surface of: a) C100 coating tested at 700 ºC, b) C140 coating

tested at 700 ºC, c) gray cast iron tested at 700 ºC.

Conclusions

Nickel based hardfaced coating deposited by plasma transferred arc on gray cast iron

with two dissimilar levels of dilution, were tribologically tested at room and high

temperatures and compared to uncoated gray iron. The results showed that the wear loss of

the coatings was independent of the substrate dilution at room temperature; tribo-oxidation

was the main wear mechanism. Despite its lower hardness, gray cast iron showed similar wear

rate as Ni-based coatings due to the lubricious effect of a C tribo-layer formed on the sliding

surfaces. Increase sliding test temperature to 550 ºC led to an abrupt increase of the wear rate

of the coatings and the gray cast iron due to the material softening and oxidation. Gray cast

iron showed extremely high wear values due to the continuous removal of the oxides, which

exposed the soft non-oxidized material. Ni-O based layer protected more efficiently the Ni-

based coatings. Further increase in test temperature to 700 ºC led to the improvement of the

wear resistance of coatings and an intensification of the wear on gray cast iron. The coating

produced with the highest dilution displayed the highest wear resistance at elevated

temperatures. This behavior was related to its lower oxidation resistance, which promoted the

easier formation of high amounts of oxide wear debris and their agglomeration into compact

oxide tribolayer acting as protection of wearing surfaces.

200 400 600 800 1000 1200 1400 1600 1800 2000

Fe3O

4

b)

c)

a)

Fe2O

3NiO

Re

lative inte

nsity

Raman shift [cm-1]

Annex E


Acknowledgments

This research is sponsored by FEDER funds through the program COMPETE – Programa

Operacional Factores de Competitividade – and by national funds through FCT – Fundação

para a Ciência e a Tecnologia –, under the projects PTDC/EME-TME/122116/2010, PEst-

C/EME/UI0285/2013 and CENTRO -07-0224 -FEDER -002001 (MT4MOBI), as well as the

grant (SFRH/BD/68740/2010).

References

[1] Cingi M, Arisoy F, Basman G, Sesen K. The effects of metallurgical structures of

different alloyed glass mold cast irons on the mold performance. Mater Lett. 2002;55:360-3.

[2] Luo L, Liu S, Li J, Yucheng W. Oxidation behavior of arc-sprayed FeMnCrAl/Cr3C2-

Ni9Al coatings deposited on low-carbon steel substrates. Surf Coat Technol. 2011;205:3411-

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[3] Vaßen R, Jarligo MO, Steinke T, Mack DE, Stöver D. Overview on advanced thermal

barrier coatings. Surf Coat Technol. 2010;205:938-42.

[4] Gatto A, Bassoli E, Fornari M. Plasma Transferred Arc deposition of powdered high

performances alloys: process parameters optimisation as a function of alloy and geometrical

configuration. Surf Coat Technol. 2004;187:265-71.

[5] Siva K, Murugan N, Logesh R. Optimization of weld bead geometry in plasma transferred

arc hardfaced austenitic stainless steel plates using genetic algorithm. Int J Adv Manuf

Technol. 2009;41:24-30.

[6] Guo C, Zhou J, Chen J, Zhao J, Yu Y, Zhou H. High temperature wear resistance of laser

cladding NiCrBSi and NiCrBSi/WC-Ni composite coatings. Wear. 2011;270:492-8.

[7] Gurumoorthy K, Kamaraj M, Rao KP, Rao AS, Venugopal S. Microstructural aspects of

plasma transferred arc surfaced Ni-based hardfacing alloy. Materials Science and

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Annex E


[8] Li W, Li Y, Sun C, Hu Z, Liang T, Lai W. Microstructural characteristics and degradation

mechanism of the NiCrAlY/CrN/DSM11 system during thermal exposure at 1100 °C. J

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[9] Balasubramanian V, Lakshminarayanan AK, Varahamoorthy R, Babu S. Application of

Response Surface Methodolody to Prediction of Dilution in Plasma Transferred Arc

Hardfacing of Stainless Steel on Carbon Steel. J Iron Steel Res Int. 2009;16:44-53.

[10] Fernandes F, Cavaleiro A, Loureiro A. Oxidation behavior of Ni-based coatings

deposited by PTA on gray cast iron. Surf Coat Technol. 2012;207:196-203.

[11] Fernandes F, Lopes B, Cavaleiro A, Ramalho A, Loureiro A. Effect of arc current on

microstructure and wear characteristics of a Ni-based coating deposited by PTA on gray cast

iron. Surf Coat Technol. 2011;205:4094-106.

[12] Fernandes F, Ramalho A, Loureiro A, Cavaleiro A. Wear resistance of a nickel-based

coating deposited by PTA on gray cast iron. Int J Surf Sci Eng. 2012;6:201-13.

[13] Vadiraj A, Balachandran G, Kamaraj M, Gopalakrishna B, Venkateshwara Rao D. Wear

behavior of alloyed hypereutectic gray cast iron. Tribology International. 2010;43:647-53.

[14] Prasad BK. Sliding wear response of a gray cast iron: Effects of some experimental

parameters. Tribology International. 2011;44:660-7.

[15] Kesavan D, Kamaraj M. The microstructure and high temperature wear performance of a

nickel base hardfaced coating. Surf Coat Technol. 2010;204:4034-43.

[16] Mironova-Ulmane N, Kuzmin A, Steins I, Grabis J, Sildos I, Pärs M. Raman scattering

in nanosized nickel oxide NiO. Journal of Physics: Conference Series. 2007;93:012039.

[17] Jian S-R, Chen Y-T, Wang C-F, Wen H-C, Chiu W-M, Yang C-S. The Influences of H2

Plasma Pretreatment on the Growth of Vertically Aligned Carbon Nanotubes by Microwave

Plasma Chemical Vapor Deposition. Nanoscale Research Letters. 2008;3:230-5.

Annex E


[18] Pandya SN, Nath SK, Chaudhary GP. Friction and Wear Characteristics of TIG

Processed Surface Modified Gray Cast Iron2009.

[19] de Faria DLA, Venâncio Silva S, de Oliveira MT. Raman microspectroscopy of some

iron oxides and oxyhydroxides. Journal of Raman Spectroscopy. 1997;28:873-8.

[20] Ferrari AC, Robertson J. Raman spectroscopy of amorphous, nanostructured, diamond–

like carbon, and nanodiamond. Philosophical Transactions of the Royal Society of London

Series A: Mathematical, Physical and Engineering Sciences. 2004;362:2477-512.

[21] Li YQ, Qiu T. Oxidation behavior of boron carbide powder. Materials Science and

Engineering: A. 2007;444:184-91.

[22] Lavrenko VA, Gogotsi YG. Influence of oxidation on the composition and structure of

the surface layer of hot-pressed boron carbide. Oxid Met. 1988;29:193-202.

Annex F


Annex F

F. Fernandes, A. Ramalho, A. Loureiro, J.M. Guilemany, M. Torrell, A.

Cavaleiro, Influence of nanostructured ZrO2 additions on the wear resistance of Ni-

based alloy coatings deposited by APS process, Wear, 303 (2013) 591-601.

Wear 303 (2013) 591–601


Wear

0043-16http://d

n CorrE-m

journal homepage: www.elsevier.com/locate/wear

Influence of nanostructured ZrO2 additions on the wear resistanceof Ni-based alloy coatings deposited by APS process

F. Fernandes a,n, A. Ramalho a, A. Loureiro a, J.M. Guilemany b, M. Torrell b, A. Cavaleiro a

a CEMUC—Department of Mechanical Engineering, University of Coimbra, Rua Luís Reis Santos, 3030-788 Coimbra, Portugalb Thermal Spray Centre (CPT-UB), University of Barcelona, C/Martí i Franques n 1, Barcelona, Spain


Article history:Received 19 December 2012Received in revised form8 April 2013Accepted 15 April 2013Available online 22 April 2013

Keywords:Sliding wearThermal spray coatingsMetal–matrix compositeSurface analysisWear testing

48/$ - see front matter & 2013 Elsevier B.V. Ax.doi.org/10.1016/j.wear.2013.04.012

esponding author. Tel.: +351 239 790 745; faxail address: [email protected] (F. Fer

a b s t r a c t

In the present investigation, the influence of the addition of nanostructure zirconia particles on themicrostructure, micro-hardness and wear performance of a Ni-based alloy (Colmonoy 88) deposited byatmospheric plasma spraying (APS) on low carbon steel has been reported. Two different procedureswere tested: (i) spraying powders of Colmonoy 88 and zirconia mixed by mechanical alloying and (ii)spraying powders separately using a dual powder injection system available at the APS equipment. Themicrostructure and the mechanical properties of coatings were characterized by scanning electronmicroscopy/energy dispersive X-ray analysis (SEM-EDS), X-ray diffraction (XRD) and micro-hardnessmeasurements. The tribological properties were evaluated at room temperature in reciprocating slidingwear equipment. The results indicate that the as-sprayed modified coatings were mainly composed byNi, Ni–Cr–Fe, Cr23C6, Cr5B3, and tetragonal zirconia. Evenly distribution of zirconia can be seen in thecoatings produced by powders prepared by mechanical alloying, while dispersive ones can be seen in theother case. Hardness and wear resistance of coatings is increased with nanostructured zirconia additions,while their friction coefficient is decreased. Coatings produced with mechanical alloying show thehighest wear resistance of all tested coatings. Nanostructured ZrO2 coating displays the worst wearresistance.

& 2013 Elsevier B.V. All rights reserved.

1. Introduction

Improvement of the surface properties by thermal sprayinghard and wear resistant materials is a commonly used industrialpractice [1–3]. The base material provides the overall mechanicalstrength of the components while coatings provide a way ofextending the limits of their use at the upper end of theirperformance capabilities.

A wide variety of coating materials such as cobalt, iron andnickel alloys can be successfully applied by thermal sprayingtechnologies to protect the surface of components subjected toharsh environments [4–6]. Superalloys and cermet APS (atmo-spheric plasma spraying) coatings have been widely employed toimprove the oxidation, corrosion, abrasion and wear resistance ofengineering components, such as plungers, molds, wearing plats,turbines, tools, etc, whose surface is submitted to extreme tribo-logical conditions in service [7–10]. Nickel-based alloys have beenespecially used in the protection of parts due to their uniquecombination of properties (mechanical, tribological, and hightemperature properties) [11–14]. The microstructure and the wear

ll rights reserved.

: +351 239 790 701.nandes).

behavior of cobalt, iron and nickel alloys coatings have beenstudied using several alloy compositions, several processes ofdeposition and different substrates. Regarding the wear resistanceof these coatings, some special composite systems were studiedand developed. The incorporation of hard and stable carbides andoxide phases, such as CrC, SiC, WC, Al2O3, TiO2, ZrO2, CeO2, etc, inthese coatings, has been reported as a way to improve their wearresistance. For example, Hou et al. [15] showed the beneficialeffect of nano-Al2O3 particles on the microstructure and wearresistance of a nickel-based alloy coating deposited by plasmatransferred arc (PTA). Harsha et al. [1] studied the influence of CrCon the microstructure, microhardness and wear resistance of anickel-based alloy deposited by flame spraying process. Theyobserved that the wear resistance of the Ni modified coating isincreased relatively to the CrC-free Ni-based alloy. Wang et al. [16]investigated the effect of three nano-particles additions (Al2O3,SiC, CeO2) on the high temperature wear behavior of a Ni-basedalloy coatings produced by laser cladding technique and reportedthat the addition of these nanoparticles increased the wearresistance of the coatings. Regarding the effect of incorporationof ZrO2 particles in a Ni matrix, it has also been extensively studied[17–20]; however, only electrodeposited coatings have been con-cerned. Despite of these studies to improve the wear resistance ofNi electrodeposited coatings, at our knowlodgment no reports



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F. Fernandes et al. / Wear 303 (2013) 591–601592

have been published as regards to the effect of nanostructuredzirconia additions on the properties of Ni-based thermal sprayedcoatings. Since zirconia is a ceramic material of high level of

Table 1Nominal chemical composition (wt.%) of the as-received powders.

Powders W B C Cr Fe Si Ni

Colmonoy 88 13.5 2.6 0.6 14.1 3.5 3.7 Balance

HfO2 Y2O3 ZrO2

Nanostructured ZrO2 2.5 7.5 Balance

Table 2Plasma spraying parameters of the dissimilar coatings.

Parameter Ni ZrO2 Ni+2

Arc current (A) 550 600 550Voltage (V) 66 74 75Gun speed (mm/s) 1250 1250 1250Leep (mm) 5 5 5Spray distance (mm) 130 60 60Primary gas (Ar) flow rate (l/min) 54 34 34Secondary gas (H2) flow rate (l/min) 9 12 15Carrier gas (O) flow rate (l/min) 2.3 3 ZrO2/Powder feed rate (rpm) 25 30 ZrO2/Agitation (rpm) 40 40 ZrO2/Injector diameter (mm) 2 2 2Number of passes 12 12 12Angle 90 90 90Substrate rotation (rpm) – – –

Injector(s) position (0) 90 90 90

Fig. 1. SEM morphologies of: (a) nickel-based alloy powders, (b) nanostructured zirconia

hardness, wear and oxidation resistance, its use as reinforcementmaterial in nickel-based alloy coatings deposited by thermalspraying processes should be investigated.

Therefore, the main goal of the present work was to study theeffect of nanostructured zirconia additions on the microstructure,micro-hardness and wear performance of a nickel-based hardfacingalloy deposited on low carbon steel by atmospheric plasma spraying(APS). The coatings with nanostructured zirconia additions wereproduced by spraying either powders prepared by mechanical alloy-ing or separate powders using a dual system of powder injection.The microstructure, the mechanical and tribological propertieswere characterized by scanning electron microscopy/energy disper-sive X-ray analysis (SEM-EDAX), X-ray diffraction (XRD), Vickersindentation tests and reciprocating sliding wear tests. All the

0% Dual Ni+40% Dual Ni+20% MA Ni+40% MA

550 550 60073 73 731250 750 7505 – –

60 120 12034 34 3415 15 15

Ni 3/2.3 ZrO2/Ni 3/2.3 3 3Ni 5/20 ZrO2/Ni 10/15 15 15Ni 40/40 ZrO2/Ni 40/40 40 40

2 2 212 28 2890 90 90– 16 1690 90 90

powders, (c) powders prepared by mechanical alloying, (condition with 40% ZrO2).

F. Fernandes et al. / Wear 303 (2013) 591–601 593

properties of the doped coatings were compared with the un-modified nickel and nanostructured zirconia coatings.


2.1. Materials

In this investigation commercial powders of a Ni-based alloy(Colmonoy 88) and nanostructured zirconia from the “Wall Colmonoy”and “Innovnano” companies, respectively, were used as feedstock toproduce coatings with Atmospheric Plasma Spraying process (APS)onto a low carbon steel substrate (AISI 8620). The Ni-based alloypowders have been characterized as a spherical morphology with anaverage size D50 of 45 mm and density of 7900 (kg/m3), whilstnanostructured ZrO2 powders have an average primary particle sizeof 60 nm with granule size D50 of 60 mm and density 1700 (kg/m3).The chemical compositions of the as-received powders are displayedin Table 1. The plasma sprayed coatings were deposited with an APSA3000 system from “Sulzer Metco”, using either pure as-receivedpowders and a dual system for powder injection available at theAPS equipment (designated as “Dual”) or mixtures of powdersprepared by mechanical alloying (MA) process (designated as “MA”).The substrate material was grit-blasted with alumina grade 24 from“Alodur Germany” before coating (which induced a correspondingsuperficial roughness (Ra) of 5 mm), to increase the surface roughnessand achieve the proper mechanical interlocking between the coatingand the substrate. For pure colmonoy and zirconia coatings and Dualprocedure, the substrates were fixed positioned, being the spraying ofthe surface ensured by the movement of the torch. On the other hand,

Fig. 2. X-ray diffraction pattern of the dissimilar powders used in the depositions.

Fig. 3. Energy dispersive spectroscopy analysis (EDS) of a powde

in the MA method the substrates were placed in a rotatory table andthe spraying was performed by combining this rotation with thevertical movement of the torch. Before final depositions, preliminarycoatings were produced changing the main deposition parametersand, then, characterized in order to properly select the depositionsconditions that allowed achieving the best results. Moreover, duringdepositions the substrates were cooled down by two pressurized airguns keeping the temperature below 150 1C, in order to avoid crackingand residual stresses formation.

Powders mixtures were prepared by mechanical alloying in aFritsch planetary ball mill using a 250ml hardened steel vial andfifteen balls with 20 mm diameter of the same material. Two propor-tions of nanostructured zirconia were added to the nickel-based alloypowder, 20% and 40% in volume, i.e. 5.1 and 12.5 gr of nanostructuredZrO2 powder for 94.9 and 87.4 gr of nickel-based alloy powder,respectively. The MA process was carried out at 400 rpm for 2 h in aprotective atmosphere of argon. After 1 h of milling, the process wasinterrupted for 10 min to cool the vial and to reserve rotation. Finally,after milling, the powder was mechanical sieved to obtain particleswith a size distribution in the range [28–71 mm]. Two differentcoatings also containing nanostructured zirconia additions of 20%and 40% were also performed by using the dual system of powderinjection available at the APS equipment. In this case the proportionwas easily achieved by controlling the flow of powders through therotation control of the powder feeders. Henceforth, coatings producedusing as received pure powder of nickel-based alloy and nanostruc-tured zirconia will be designated as “Ni”, and “ZrO2”, respectively,coatings produced with powders prepared by mechanical alloyingusing 20% and 40% of nanostructured ZrO2 will be termed as “Ni+20%MA” and “Ni+40% MA”, respectively, and coatings deposited using thedual system of powder injection with 20% and 40% of nanostructuredZrO2 will be designated as “Ni+20% Dual” and “Ni+40% Dual”,respectively. The process parameters employed in the dissimilarcoating depositions are shown in Table 2.

2.2. Powders and coatings characterization

Powders morphology were observed and characterized by scan-ning electron microscopy with X-ray spectroscopy (SEM/EDS) andX-ray diffraction (XRD) using Co Kα radiation. Transverse sectionof APS coatings was polished using conventional metallographicprocedure which consisted grinding followed by polishing. Themicrostructure of coatings was studied and characterized by opticalmicroscope (OM) and SEM/EDS. Further, the structure of the coatingswas studied by X-ray diffraction. The micro-hardness of coatings wasdetermined by Vickers testing using a 3 N load at 15 s holding time. Atotal of 30 measurements of hardness were done in each coating.

r produced by mechanical alloying (condition Ni+40% ZrO2).


2.3. Wear tests

Wear behavior of APS coatings was studied using reciprocatingwear equipment. An harmonic wave generated by an eccentric androd mechanism imposed a stroke length of 2.05 mm at a frequencyof 1 Hz. A detailed description of the equipment can be found inreference [21]. Before wear tests, the surface of coatings weregrounded and then polished using a 3 mmdiamond past. A superficialroughness Ra of: 0.51, 0.65, 0.49 and 0.98 mm was measuredrespectively for Ni, Ni Dual, Ni MA and ZrO2 coatings. In the scopeof the present study, the values of normal load applied to the coatedsamples were 5, 7, 8.5 and 10 N. All the tests were conducted at aconstant rotational speed of 190 rpm during 2 h. The acquisition time

Fig. 4. SEM morphology of the: (a) Ni-based alloy coating, (b) nanostructured ZrO2 coatand Ni+40% MA.

of signal used was 60 s. A soda-lime glass sphere with 500 HV0.5 ofhardness and 10 mm in diameter was used as counterpart, in orderto study the interaction of glass with the different coatings. Afterwear tests, the volume loss was estimated through the transverseprofile of each wear scar at the middle and near the end of both sidesof the scar. To ensure the reproducibility of results, a set of three scarsfor each test condition was performed at the coating, and the volumeloss calculated by using the volume average of these marks. Anapproach of the Archard's law was applied; see Eq. (1), in order toobtain the specific wear rate. In this equation, V represents the wearvolume, P is the normal load, k is the specific wear rate and l thesliding distance. To estimate the error of the measurements, astatistical analysis described elsewhere by A. Ramalho [22] was used.

ing, (c) and (d) Ni+20% Dual and Ni+40% Dual, respectively, (e) and (f) Ni+20% MA


After wear tests, the morphology of depressions was observed andcharacterized by scanning electron microscopy.

V ¼ k� P � l ð1Þ


3.1. Powders characterization

Fig. 1(a–c) shows SEM morphologies of nickel-based alloy,nanostructured zirconia and one of the MA prepared powders.

Fig. 5. SEM-EDS analysis performed at the: (a) Ni coating, (b) Ni+40% Dual coating, (c) N

Ni and ZrO2 powders clearly display spherical-shaped particles, whileMA powders have irregular shapes. XRD patterns in Fig. 2 shows thatNi powders consist essentially in a Ni solid solution (ICDD card 01-1258) with some traces of chromium carbide (ICDD card 85-1281)and tungsten (ICDD card 47-1319). This indexation is in agreementwith the results from the literature where similar phases weredetected [23]. Nanostructured ZrO2 powders are essentially formedby tetragonal zirconia. In respect to MA powders XRD pattern revealthe same phases indexed from the pure Ni and nanostructured ZrO2

powders, suggesting only physical mixing, not having occurredchemical reactions. Fig. 3 shows a SEM-EDS analysis performed at a

i+40% MA coating. SEM-EDS spectra of: (d) point 1, (e) zone 2, (f) point 3, (g) zone 4.

Fig. 6. X-ray diffraction pattern of the dissimilar coatings.

Fig. 7. Microhardness of coatings.

Fig. 8. Volume loss of: (a) coatings, (b) ball; after wear tests as function of theapplied load.

Fig. 9. Evolution of the wear behavior as a function of the parameter “normal load� sliding distance”.

Table 3Results of the linearization analysis of the dissimilar coatings.

CoatingSpecific wear rate, k(mm3/N�m)

Average STD confidence interval 90% r2

Ni 1.68�10−5 3.06�10−6 7.83�10−6–2.57�10−5 0.938Ni+20% Dual 4.45�10−6 1.14�10−6 1.12�10−6–7.79�10−6 0.884Ni+40% Dual 5.26�10−6 8.65�10−7 2.74�10−6–7.78�10−6 0.949Ni+20% MA 3.83�10−6 5.56�10−7 2.21�10−6–5.46�10−6 0.960Ni+40% MA 1.90�10−6 3.97�10−7 7.38�10−7–3.06�10−6 0.919ZrO2 3.03�10−4 7.15�10−5 9.48�10−5–5.12�10−4 0.900


MA powder (Ni+40% ZrO2). As the EDS spectra displays, the powderis essentially formed by nickel, zirconia, oxygen and tungsten,supporting the results from the XRD analysis.

3.2. Microstructure

Fig. 4 shows SEM micrographs of the cross section of thecoatings. All coatings present the typical morphology of plasmaspraying, with pores, lamellae, partially melted and un-meltedparticles and good adhesion to the substrate. ZrO2 (see Fig. 4b) andNi+ZrO2 Dual coatings (Fig. 4(c and d)) show a higher level ofporosity than the other coatings. Very often, small pores werefound within the flattened particles and result from shrinkageporosity and big pores were normally located between lamellaeand were caused by gas porosity phenomenon [24]. In all cases themelting states of particles influence the level of porosity and,usually, low melting state of particles gives rise to a porousmicrostructure. Furthermore, pure ZrO2 coating, (see Fig. 4b)reveals the presence of some cracks in the coating, resulting fromtensile residual stresses generated during cooling down to roomtemperature. Ni and MA coatings (see Fig. 4(a), (e) and (f)) displaya compact and homogenous structure, while Dual coatingsrevealed a microstructure full of semi-melted particles, suggestinga brittle character based on low level of agglomeration betweenparticles. It is well known that in the thermal spraying processesthe melting level depends on the thermal energy added to theparticles, on the plasma temperature and on the particles flyingtime. During plasma spraying, molten particles impact the surface


of substrate, then deform, solidify and transform into lamellae.During spraying, the particles can be in all of the followingstates on impact: fully molten, superheated, semi-molten andmolten then re-solidified [24,25], depending on the depositionparameters. It is well known that in plasma spraying the electricarc current, the primary plasma gas flow rate, the second plasmagas flow rate and powder size are the main parameters thatinfluence the in-flight particle behaviors. From Table 2, comparingthe process parameters for the Ni and Ni+20% and 40% ZrO2 Dualcoatings, differences can be observed in relation to the spraydistance, primary and secondary gas flow rates, and powder feedrate. According to the literature, an increase of argon flow rateleads to a gentle decrease of particle temperature but a rapidincrease of velocity [26]. On the other hand, increasing thehydrogen flow rate increases both, temperature and velocity,though temperature is more sensitive to the hydrogen flow ratethan velocity. So, due to the lower primary gas flow rate (Ar) andhigher secondary gas flow rate (H2) used in the Ni +ZrO2 Dualcoatings, would be expected a great level of melting particles andconsequently an homogenous structure, if the particles do notoxidize during the in-flight time. However, such fact was notobserved. The lower primary gas flow rate (Ar) used in theproduction of these coatings might suggest that independentlyof the high melting state of particles, the velocity given by the flowcould be not enough to accelerate the particles onto the specimensurface. Nevertheless, the same process parameters were used inthe deposition of the Ni+20% MA coating, with exception of thespraying distance, and the microstructure revealed to be homo-genous and compact. In fact, the spraying distance also plays animportant role in the thermal treatment of particles during theirin-flight time [27]. Thereby, for short spraying distances, due tothe short exposure of particles in the flame, the energy added bythe flame could be not enough to semi-molten or molten theparticles, giving rise to a microstructure similar to those observedfor these coatings. Furthermore, it can also be speculated that thetwo different ways to add zirconia at the nickel alloy can changethe entropy of the system, and therefore change the running-on of

Fig. 10. Variation of the friction force of coatings with applied load for: (a) Ni coating,

the deposition. In fact, according to the literature the nature ofpowders has a detrimental effect in their interaction with theplasma spraying flame and, therefore, on the properties of thecoatings [28]. The low level of melting particles can also explainthe high level of porosity observed in these coatings.

For better characterization and perception of the differentcoatings microstructure, magnifications of the Ni, Ni+40% Dualand Ni+40% MA coatings, as well as EDS analyses are shown inFig. 5. The EDS analysis revealed that Ni coating is essentiallyformed by Ni, W, Cr and Fe, which matches well with the chemicalcomposition of the as received Ni powders. Moreover, smallparticles rich in W (white phase) can be observed evenly dis-tributed in the microstructure. The semi-melted particles in themicrostructure of the Ni+40% Dual coating are from the samenature of the Ni coating matrix, therefore, they can be identified assemi-melted Colmonoy powders. Similar to Ni coating a whitephase rich in W is also detected in the microstructure, co-existingwith a coarse gray phase, entrapped between the boundaries ofthe semi-melted particles, revealed by EDS to be zirconia. Such adistribution is in the basis of a brittle behavior. Contrary to thiscoating, Ni+40% MA displays a compact and homogeneous micro-structure, with small zirconia particles evenly distributed alongthe microstructure improving the coating toughness. Therefore,from these results, it can be inferred that the addition of nano-structured zirconia can be successfully achieved by using bothprocedures, powders prepared by mechanical alloying or dualprojection. However, in coatings deposited with MA powderszirconia is finer and more homogeneously distributed in theNi-matrix giving rise to better mechanical performance.

XRD patterns of the coatings are displayed in Fig. 6. As thespectrum reveals, the nickel-based coating consists mainly of Ni,Ni-Cr-Fe, Cr23C6 and Cr5B3 phases. These results are in accordancewith the literature where similar phases were identified [23].Moreover, all the coatings with nanostructured zirconia additionsshowed diffraction lines corresponding to the same phases identi-fied at both the Ni and ZrO2 coatings. In the latter, is mainlyformed by tetragonal zirconia.

(b) Ni+40% Dual coating, (c) Ni+40% MA coating (d) nanostructured ZrO2 coating.


3.3. Micro-hardness

Fig. 7 shows that nanostructured ZrO2 additions increases thehardness of the nickel-based alloy. The same behavior was reportedby other authors after addition of hard particles such as (Al2O3, CeO2,SiC, WC) at a nickel alloy [1,3,16]. However, MA coatings have higherhardness than Dual ones. This result is coherent with either thelower porosity or the finer and more homogeneous distribution ofnanostructured zirconia in Ni+ZrO2 MA coatings.

3.4. Wear behavior

The volume loss due to wear of the coatings and counterpartincreases with increasing applied load, as it is shown in Fig. 8(a) and(b), respectively, being this trendmore notorious for ZrO2 coating. As itwill be shown later, the higher volume loss for this coating can becorrelated with the brittle behavior as displayed on its worn surface.Nanostructured zirconia additions have a positive effect, decreasingthe volume loss of material, independently of the deposition proce-dure used. However, efficiency seems to be higher in Ni+ZrO2 MAcoatings which display higher decreases of the volume loss than Ni+ZrO2 Dual coatings. Furthermore, increasing zirconia content in MAcoatings continues to have a beneficial effect whereas the inverse isobserved in Dual coatings. The better wear performance of Ni+ZrO2

MA coatings can be attributed to the combination of a higherhardness, an evenly distribution of smaller zirconia domains in theNi-based matrix and a more compact microstructure. In fact, thepresence of hard phases increases the overall hardness of the coatingand, on the other hand, smaller and uniformly distributed hard

Fig. 11. SEM morphologies of the worn surfaces tested at 8.5 N load of the:

domains in a soft matrix gives rise to tougher materials, whichcombined lead to a decrease of the volume loss due to wear. Similarbehavior was observed by other author after adding hard particles to anickel based alloy [1,15]. In the case of Ni+ZrO2 Dual coatingsnanostructured hard ZrO2 phases have initially a similar influence,by increasing the hardness. However, the agglomeration of ZrO2

particles, localized in-between Ni-based lamellae makes microstruc-ture less tough with its consequent detrimental effect on the wearperformance when high contents of the hard phase exist.

In order to normalize the wear results, they were plotted in Fig. 9in terms of material volume loss as function of the product of slidingdistance by the applied load. For each set of data, a straight lineartrend was fitted in order to obtain the specific wear rate from theslope (see Eq. 1). Further, a complete reliability analysis of the specificwear rate is presented in Table 3 for a confidence interval of 90%.Because of the high volume loss displayed by the ZrO2 coating, onlythe point produced with lower load and the beginning of the straightadjustment were plotted in Fig. 9. Crossing the information of Fig. 9and Table 3 it is possible conclude the specific wear rate ofNi-modified coatings is lower than that of the un-modified Ni coating.Ni+40% ZrO2 dual coating display the lowest specific wear rate, beingthis result attributed to the higher hardness displayed by this coating.Further, coatings produced by the dual system of powder injectiondisplay higher specific wear rate than coatings produced by powdersprepared by mechanical alloying. ZrO2 coating displays the highestspecific wear rate, being one and two orders of magnitude greaterthan the un-modified and modified coatings, respectively.

The average values of the frictional force of the Ni, Ni+40% Dual,Ni+40% MA and ZrO2 coatings linearly increases with the applied

(a) Ni coating, (b) Ni+40% dual, (c) Ni+40% MA, (d) ZrO2 coating.


load, as shown in Fig. 10. This is in good agreement with Amontons–Coulomb model. By fitting a linear trend to the experimentalresults, the friction coefficient calculated from the slope decreases

Fig. 12. Magnification of the worn surfaces tested with a load of 8.5 N of the: (a) Ni(e) point A, (f) point B.

with the nanostructured zirconia additions, independently of thedeposition method. Pure nanostructured zirconia shows higherfriction coefficient than composite coatings, but lower than the

coating, (b) Ni+40% dual, (c) Ni+40% MA, (d) ZrO2 coating. SEM-EDS spectra of:


pure nickel one. The changes in the specific wear rate and frictioncoefficient described above are only partially correlated with theapplied test conditions and properties of the materials, which canproduce dissimilar wear mechanisms during wear tests, and there-fore dissimilar volume loss of coatings. These reasons have encour-aged the investigation of the interaction between the pairspecimens-counterpart. Scanning electron microscopy (SEM) ana-lysis was conducted on the dissimilar worn surfaces of coatings inorder to identify the main wear mechanisms produced.

Typical SEM morphologies of worn surfaces of the coatingstested with 8.5 N against a glass sphere are shown in Fig. 11.Distinct wear mechanisms were identified in the worn surfaces. InNi and Ni+ZrO2 Dual coatings significant amount of adherentmaterial is shown in the wear scar, while Ni+ZrO2 MA and ZrO2

coatings show clean worn surfaces. EDS analysis performed inadhered debris, marked in Fig. 12(a) and (b) by letters A and B,allows concluding that they are made of silicon oxides and Ni,suggesting a mechanical mixing by the friction process, joining thesofter Ni phase with the silica type wear debris from the ball wear.Being softer, the compact wear debris can stick at the worn trackas well as at the ball wear scar, as can be observed in Fig. 13.During contact the friction generated by the movement of thecounterpart against the surface specimen, causes an increase oftemperature promoting local plastic deformation. At this time,both materials will adhere to each other. In the wear track of theNi coating, see Fig. 12(a), the wear debris tend to be incorporatedin the specimens, more precisely in the boundaries of the semi-melted powders, remaining adherent and leading to removal ofmaterial. This mechanism explains simultaneously the high fric-tion values reported for Ni coating. On the other hand, ZrO2

coating shows clean wear scars without traces of adhered material.The surface has a brittle appearance, where micro cracks can beseen, as shown in Fig. 12(d), revealing a rough surface from whichparticles of several tens of micrometers were detached. Thismorphology suggests mechanical interlocking during the contactwhich can explain the higher values of the friction coefficient.Furthermore, it should promote a strong abrasion in the counter-part, giving rise to much higher wear rate as shown in Fig. 8(b).Moreover, the evolution of the wear amount with the increase ofnormal load, (see Fig. 8), allow identifying a made severe wearvolume for loads higher than 7 N, which should be correlated to atransition for a made brittle behavior. Similar behavior wasobserved by Chen et al. [29] during the evaluation ofun-lubricated wear properties of plasma-sprayed nanostructuresand conventional zirconia coatings. The wear scars of Ni+ZrO2

Fig. 13. Optical micrograph of the ball wear scars tested against the Ni coating with7 N of load.

Dual coatings show a mix morphology of pure Ni and ZrO2

coatings, i.e. adhered debris can be observed over a roughbrittle-appearance surface, whose number decreases with nanos-tructured zirconia additions. In these coatings, zirconia particlesare localized in the boundaries of the semi-melted Ni-basedpowders (see Fig. 4c and d). Despite of the particles operating asreinforcement of the coatings, leading to a lower volume loss, thearea of the nickel exposed to the glass sphere remains higher.Therefore, it is expected similar debris adhesion as for Ni coating.At the same time, due to the continuous removal of material, ZrO2

particles in the boundaries of the semi-melted powders loses theirsupport capacity being removed from the boundaries. Theseparticles will slide through the wear track acting as a removal ofthe wear debris. Increasing the volume of nanostructured zirconiafrom 20 to 40%, higher levels of brittleness in the boundaries areachieved, increasing the volume loss of material, as reportedbefore, and giving rise to the rough and brittle appearance shownbelow the adhered debris. The higher level of porosity of thesecoatings also accounts for higher levels of volume loss of material.

The coatings produced from MA powders show also clean wearsurfaces but with a different failure type from pure zirconia coating. Ni+ZrO2 MA coatings reveal surfaces with scratches identifying somesort of grooving abrasion, see Fig. 12(c). The uniform distribution ofsmall nanostructured zirconia particles and the higher hardness andtoughness of these coatings avoids the plastic deformation, theadhesion and the liberation of large wear debris either from thecoating or the counterpart. The smooth worn surfaces can be in thebasis of the low friction coefficient [30]. In these coatings, the matrixcutting occurs through a process of two body abrasion, as shownFig. 12(c). Despite of the dissimilar wear mechanisms observed in thewear track of the coatings, during the wear tests it was observed thatmost of the material resulting from the wear was blown out, due tothe movement of the ball.

4. Conclusions

�
The influence of nanostructured zirconia additions on themicrostructure, micro-hardness and wear performance of anickel-based hardfacing alloy deposited by atmosphericplasma spraying (APS) on low carbon steel is reported in thepresent paper.
�
The coatings containing nanostructured zirconia additionswere produced either using powders prepared by mechanicalalloying or using separate powders that are sprayed simulta-neously via a dual system of powder injection.
�
Nanostructured zirconia can be successfully incorporated intonickel-based alloy powders by mechanical alloying. Coatingsproduced with these powders present evenly distributed zir-conia as opposed to coatings produced by separate powderinjection, where the nanostructured zirconia particles wereentrapped in the boundaries of the semi-melted Ni lamellae.
�
The microstructure of the nickel coating consists mainly of Ni,Ni–Cr–Fe, Cr23C6 and Cr5B3 phases. The same diffraction lineswere identified in the modified nickel coatings with tetragonalzirconia, independently of the procedure used in depositions.
�
The hardness and wear behavior of coatings was improvedwith nanostructured zirconia additions while the frictioncoefficient is decreased.
�
The nanostructured ZrO2 coating shows the highest hardnessvalues but also the highest specific wear rate, one to two ordersof magnitude higher than the other prepared coatings, inparticular those deposited with MA powders.
�
Adhesive wear mechanism was observed in the worn surface ofthe nickel coating and coatings prepared by Dual deposition,


while abrasive wear was observed in ZrO2 coating and coatingsdeposited using MA powders.

Acknowledgments

The authors wish to express their sincere thanks to thePortuguese Foundation for the Science and Technology (FCT),through COMPETE program from QREN and to FEDER, for financialsupport in the aim of the Project number “13545”, as well as for theGrant (SFRH/BD/68740/2010).

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Annex G


Annex G

F. Fernandes, A. Loureiro, T. Polcar, A. Cavaleiro, "The effect of increasing V

content on the structure, mechanical properties and oxidation resistance of Ti-Si-V-

N films deposited by DC reactive magnetron sputtering, Applied Surface Science,

289, (2014) 114-123.

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Contents lists available at ScienceDirect

Applied Surface Science

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he effect of increasing V content on the structure, mechanicalroperties and oxidation resistance of Ti–Si–V–N films deposited byC reactive magnetron sputtering

. Fernandesa,∗, A. Loureiroa, T. Polcarb,c, A. Cavaleiroa

CEMUC—Department of Mechanical Engineering, University of Coimbra, Rua Luís Reis Santos, 3030-788 Coimbra, PortugalDepartment of Control Engineering, Czech Technical University in Prague, Technicka 2, Prague 6 166 27, Czech RepublicnCATS, University of Southampton, Highfield Campus, SO17 1BJ Southampton, UK

r t i c l e i n f o

rticle history:eceived 26 June 2013eceived in revised form 18 October 2013ccepted 19 October 2013vailable online 28 October 2013

eywords:iSi(V)N filmstructureechanical properties

a b s t r a c t

In the last years, vanadium rich films have been introduced as possible candidates for self-lubricationat high temperatures, based on the formation of V2O5 oxide. The aim of this investigation was to studythe effect of V additions on the structure, mechanical properties and oxidation resistance of Ti–Si–V–Ncoatings deposited by DC reactive magnetron sputtering. The results achieved for TiSiVN films were com-pared and discussed in relation to TiN and TiSiN films prepared as reference. All coatings presented a fccNaCl-type structure. A shift of the diffraction peaks to higher angles with increasing Si and V contentssuggested the formation of a substitutional solid solution in TiN phase. Hardness and Young’s modulusof the coatings were similar regardless on V content. The onset of oxidation of the films decreased sig-nificantly to 500 ◦C when V was added into the films; this behaviour was independent of the Si and V

xidation resistanceanadium oxide

contents. The thermogravimetric isothermal curves of TiSiVN coatings oxidized at temperatures belowthe melting point of �-V2O5 (∼685 ◦C) showed two stages: at an early stage, the weight increase overtime is linear, whilst, in the second stage, a parabolic evolution can be fitted to the experimental data. Athigher temperatures only a parabolic evolution was fitted. �-V2O5 was the main phase detected at theoxidized surface of the coatings. Reduction of �-V2O5 to �-V2O5 phase occurred for temperatures above
its melting point.
. Introduction

High-speed cutting and dry machining processes without these of environmental harmful lubrication requires milling toolsapable to withstand severe conditions of wear, friction, oxidationnd corrosion [1,2]. Solid lubricants coatings, such as WC/C, MoS2,iamond-like carbon (DLC), hex-BN as well as their combinations

n nanocrystaline or multilayer structures, have been successfullypplied in such parts in order to increase their lifetime and perfor-ance for various tribological applications. However, considerable

egradation of the tribological effectiveness of these coatings at ele-ated temperature has been reported due to their low resistance toxidation [3–5]. To overcome this shortcoming, a new concept ofubrication based on the formation of lubricious oxides has been
roposed. This is the case of the so-called Magnéli oxide phasesased on Ti, Si, Mo, W and V, which have easy shear able planes3,6]. Among these elements, particular attention has been given
∗ Corresponding author. Tel.: +351 239 790 745/+351 963 239 877;ax: +351 239 790 701.

E-mail address: [email protected] (F. Fernandes).

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to the vanadium-containing coatings (Magnéli phases VnO3n−1),which showed interesting tribological properties in the temper-ature range 500–700 ◦C [3,7–12]. Dissimilar series of V-based hardcoatings have been developed, such as ternary CrVN [13], (V,Ti)N[14], multilayer AlN/VN [15] and quaternary single layer or mul-tilayered CrAlVN [16,17] and TiAlVN [2,8,18,19]. However, sinceother ternary coatings systems, based on (TiX)N with X = B, Cr, Al,Si, Cr, etc, have been attractive for advanced hard coating materi-als, the effect of vanadium doping to these systems should also beconsidered.

Among those ternary systems, most of the research works havebeen focused on TiSiN coatings deposited by CVD and/or PVD tech-niques [20,21]. Depending on the deposition conditions, these coat-ings could be deposited with a nanocomposite structure, consistingof nano-sized TiN crystallites surrounded by an amorphous matrixof Si3N4 displaying very high hardness values [22], compared to anyother ternary (TiX)N compound, with extremely good oxidationresistance [23,24]. In other deposition conditions, substitutional
solid solution of Si in TiN structure (not predicted by the Ti–Si–Nphase diagram) could be achieved [22,25,26]. Relatively to thefriction coefficient of Ti–Si–N coatings, the studies have reportedvalues in the range [0.6–1.3] at room and high temperatures
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F. Fernandes et al. / Applied Surface Science 289 (2014) 114– 123 115

Table 1Sample designation and deposition parameters of the coatings.

Sample Target power density (W/cm2) Deposition time (min) Pellets of V at the Ti target

Ti TiSi2

Interlayer Ti 10 – 5 –Interlayer TiN – 15 –

TiSiN S1-0 6.5 1 210 0S1-4 4

TiSiVN S1-8 8S1-12 12

TiSiN S2-0 6 1.5 193 0S2-4 4

TiSiVN S2-8 8S2-12 12

TiSiN S3-0 5 2.5 153 0S3-4 4

TiSiVN S3-8 8

[cooe

vmiwsa

2

situ1ecctotwswdtStTr(u1

c(las

S3-12

TiN 10 –

27–29]. Since the addition of V successfully decreased the frictionoefficient of binary and ternary systems (down to reported valuesf 0.2–0.3 at temperatures between 550 and 700 ◦C), similar studiesn the effect of V-addition should be performed on TiSiN coatingsxhibiting unique mechanical properties and oxidation resistance.

This paper reports the first results on the effect of increasinganadium content to Ti–Si–V–N films deposited by DC reactiveagnetron sputtering. It is focused on coating structure, mechan-

cal properties and oxidation resistance. Comparison of the resultsith those achieved for reference TiN and TiSiN coatings is pre-

ented. The results of this study will support further research aimedt the tribological behaviour of the coatings at high temperatures.

. Experimental procedure

Three series of TiSiN films with different Si contents, eacheries with increasing V contents (TiSiVN coatings), were depositedn a d.c. reactive magnetron sputtering machine equipped withwo rectangular (100 × 200 mm) magnetron cathodes working innbalanced mode. A high purity Ti (99.9%) target, with 18 holes of0 mm in diameter (uniformly distributed throughout the prefer-ntial erosion zone of the target), and a high purity TiSi2 (99.9%)omposite target were used in the depositions. The different sili-on contents were achieved by changing the power density appliedo each target. The V content was varied by changing the numberf high purity rods of Ti and V (with 10 mm in diameter) placed inhe holes of the Ti target. 4, 8 and 12 cylindrical pieces of vanadiumere used. In all the cases the total power applied to the targets was

et to 1500 W. To serve as reference, a stoichiometric TiN coatingas deposited from the Ti target. Hereinafter the coatings will beesignated as Sx-y, where x is related to the specific power appliedo the TiSi2 target (see Table 1), and therefore associated to thei content on the coatings, and y the number of V rods used inhe depositions, giving an indication of V content in each series ofiSiVN films. Thus, denomination S2-0, S2-4, S2-8 and S2-12 rep-esents coatings from the same series, i.e. with identical Si contentTiSi2 target power 1.5 W/cm2), with increasing V content from 0p to a maximum content achieved for the coating produced using2 vanadium pellets embedded into the Ti target.

Polished high-speed steel (AISI M2) (Ø 20 × 3 mm, for mechani-al properties measurements), FeCrAl alloy and alumina substrates
10 × 10 × 1 mm, for oxidation tests and structural analysis), stain-ess steel discs (Ø 20 × 1 mm, for residual stress measurements)nd (1 1 1) silicon samples (10 × 10 × 0.8 mm, for thickness mea-urements and chemical composition evaluation) were used as
1273 –

substrates. Prior to the depositions, all the substrates were ultra-sonically cleaned in acetone for 15 min and alcohol for 10 min. Thesubstrates were mounted in a substrate holder (which revolvedwith 18 rev/min around the centre axis) giving a target to sub-strate distance of 175 mm. Prior to deposition, the chamber wasevacuated down to 8.7 × 10−4 Pa and the substrates were etchedwith Ar ion sputtering during 1 h with a bias voltage of −650 V toremove any surface contaminants. In order to enhance the adhesionof the coatings, Ti and TiN adhesion layers of approximately 0.24and 0.45 �m, respectively, were deposited on the substrates beforeTiSi(V)N coatings. In all the depositions, the total working gas pres-sure was kept constant at 0.3 Pa, using approximately 30 sccm ofAr and 17 sccm of N2. The depositions were performed with a neg-ative substrate bias of 50 V. The deposition time was set in order toobtain films with approximately 2.5 �m of total thickness (includ-ing interlayers). The deposition parameters are shown in Table 1.

The chemical composition of the coatings was evaluated byelectron probe microanalysis (EPMA—Cameca SX 50). Crystallo-graphic structure was investigated by X-ray diffraction (X’ PertPro MPD diffractometer) using a grazing incidence angle of 1◦ andCu K�1 radiation (� = 1.54060 A). The XRD spectra were fitted byusing a pseudo-Voigt function to calculate either the full widthat half maximum (FWHM) and the peak position (2�). The frac-ture cross-section morphology and the thickness of the films wereinvestigated by scanning electron microscopy (SEM). Further, pro-filometer was used to confirm coating thickness and to evaluatesurface roughness.

The hardness and Young’s modulus of films were measured ina nano-indentation equipment (Micro Materials NanoTest) using aBerkovich diamond pyramid indenter. In order to avoid the effectof the substrate, the applied load (10 mN) was selected to keepthe indentation depth less than 10% of the coating’ thickness. 32measurements were performed in each sample. The level of resid-ual stresses was calculated through the Stoney equation [30] bymeasuring the bulk deflection of the film-substrate body.

The oxidation resistance of the coatings was evaluated by ther-mogravimetric analysis (TGA) using industrial air (99.99% purity).The films deposited on alumina substrates were firstly heated witha constant temperature ramp of 20 ◦C/min from room temperatureup to 1200 ◦C, in order to determine the onset point of oxidation.Then, coated specimens of FeCrAl alloy were subjected to isother-
mal tests at different selected temperatures and time. The weightgain of the samples was evaluated at regular 2 s intervals using amicrobalance with an accuracy of 0.01 mg. The air flux used was50 ml/min and the heating rate up to the isothermal temperature

116 F. Fernandes et al. / Applied Surface

F(

wtwst

3

3

3

bnpcro

ctSadcrtco

ig. 1. XRD patterns of: (a) TiSiN coatings with Si content ranging from 0 to 12.6 at.%,b) coating S1-0 with V additions.

as 20 ◦C/min. After oxidation tests, the surface morphology ofhe specimens was examined by scanning electron microscopyith energy dispersive X-ray spectroscopy (SEM–EDS) and their

tructure was evaluated by XRD diffraction. The oxide products onhe surface were characterized by Raman spectroscopy.

. Results and discussion

.1. Coatings characterization

.1.1. Chemical composition and microstructureThe elemental chemical composition of the coatings determined

y EPMA is shown in Table 2. As could be expected, the simulta-eous decrease in the power density of Ti target and increase in theower density of TiSi2 target gave rise to higher Si content in theoatings. Only insignificant variations were detected on the Si/Tiatios with incorporation of V in the Ti target. The nitrogen contentf the coatings was close to 50 at.% regardless of Si and V.

The corresponding X-ray diffraction patterns of the as-depositedoatings are displayed in Fig. 1. The influence of increasing V con-ent on TiSiN coatings is shown in Fig. 1b) for the film with loweri content, as representative of all series containing vanadium. Inll graphs TiN film pattern is shown as a reference. In all cases, theiffraction peaks could be generally assigned to the fcc NaCl-typerystalline TiN phase. The XRD pattern of the TiN reference coating
eveals that all peaks are shifted to lower diffraction angles in rela-ion to the standard ICDD card of TiN, probably indicating residualompressive stresses [31]. For TiSiN coatings with 3.8 and 6.7 at.%f Si, (2 0 0) peak vanishes and the sharpening of the peaks suggests
Science 289 (2014) 114– 123

larger grain size. Based on Scherrer’ equation, grain sizes of 16 and24 nm were calculated for TiN and coatings with low Si content,respectively. With further increase of the Si content (12.6 at.% ofSi) a strong decrease of the peaks intensity, as well as their signifi-cant broadening (a grain size of 6 nm was calculated), are observedsuggesting a decrease of crystallinity. With increasing Si content,the fcc peaks are shifted to higher angles, behaviour associated toa smaller unit cell. This can be explained by the presence of Si insolid solution, since its smaller atomic radius promotes the con-traction of the TiN lattice. Therefore, in this case the stoichiometricphases (TiN + Si3N4) do not segregate and nanocomposite structureis not formed. Such behaviour could be explained by the low depo-sition temperature together with low substrate ion bombardment,which do not provide the necessary mobility of the species for thenanocomposite formation [22]. This observation is in agreementwith the results from the literature, where similar solid solutionsand changes in the film orientation have been observed with siliconincorporation [21,25,32].

The coatings with V displayed similar XRD patterns and peaksto TiSiN. No significant changes in the peaks intensities or widthswere observed, except for the highest Si content films where asmall improvement in the crystallinity was detected. All the vana-dium containing films had the XRD peaks shifted to higher angles,behaviour, once again, related to a smaller unit cell due to the sub-stitution of Ti by the smaller V atoms. Similar trend was observedby Pfeiler et al. [2] for the TiAlVN system.

Fig. 2 displays typical surface and fracture cross section-micrographs of TiSiVN coatings. All coatings showed a typicalcolumnar structure. As Si content was increased (see Fig. 2b)) forcoating S1-0), the columns size and the surface roughness werefirstly reduced, for silicon contents of 3.8 and 6.7 at.%, and thenincreased for the highest silicon content, which is an oppositetrend than that observed for the grain size. The addition of V didnot change the global type of columnar morphology and the coat-ing cross-section was similar. The only exception was the strongincrease of the size of the columns as well as the surface roughness(see Fig. 2c) for coating S1-12) for low silicon content coatings.

3.1.2. Hardness, Young’s modulus and residual stressesThe dependence of the hardness and Young’ modulus on the

Si and V content is shown in Fig. 3. As the silicon content in theTiSiN films increases, the hardness of coatings reached a maximumvalue of 27 GPa at a Si content of 6.7 at%; then it dropped with fur-ther increasing Si content to a lower value than the reference TiN.Although the grain grow is observed with increasing Si additionsup to 6.7 at% of Si, the hardening of coatings is mostly result of thenitride lattice distortion, which improves the resistance to plasticdeformation (solid solution hardening). Residual stresses shouldnot be a factor of hardness differentiation between the coatings,since all films have approximately the same level of compressiveresidual stresses, with any Si addition (∼3 GPa). This result also sup-ports what was stated above that the shift of the peaks to the rightwith Si additions is exclusively due to substitutional solid solu-tion. The decrease in hardness with further Si addition is due tothe loss of crystallinity associated with the lower degree of phasenitriding regardless on the lower grain size [26]. Young’s modu-lus progressively decreases with increasing Si content, which isrelated to changes in the binding energy between ions due to Siincorporations. Similar hardness and Young’s modulus evolutionas a function of Si content has been reported in Refs [25,33].

Concerning the effect of V addition on the hardness and Young’modulus, similar evolution was observed for all films (see e.g.
Fig. 3b) for S1-y films). Although the variations are very small andoften within the error bars of the hardness values, general trend isa slight increase in hardness followed by drop for the highest vana-dium content. Coatings with the higher silicon content and with


Table 2Chemical composition of the dissimilar coatings in at.%.

Coating at.%

N O Si Ti V

TiN 49.8 ± 0.1 0.4 ± 0.1 – 49.6 ± 0.4 –

TiSiN S1-0 51.1 ± 0.6 0.4 ± 0.1 3.8 ± 0.0 44.7 ± 0.7 –S1-4 51.4 ± 0.4 1.5 ± 0.3 3.1 ± 0.0 43.0 ± 0.7 1.1 ± 0.1

TiSiVN S1-8 51.5 ± 0.2 1.2 ± 0.1 2.9 ± 0.0 37.1 ± 0.3 7.3 ± 0.1S1-12 49.7 ± 0.3 1.8 ± 0.1 2.8 ± 0.0 33.8 ± 0.3 12.0 ± 0.1

TiSiN S2-0 51.5 ± 0.2 0.6 ± 0.1 6.7 ± 0.0 41.3 ± 0.3 –S2-4 52.1 ± 0.2 1.2 ± 0.1 5.7 ± 0.0 39.3 ± 0.4 1.6 ± 0.2

TiSiVN S2-8 51.7 ± 0.1 1.4 ± 0.1 5.6 ± 0.1 33.6 ± 0.3 7.6 ± 0.2S2-12 51.1 ± 0.6 1.9 ± 0.2 5.3 ± 0.0 30.5 ± 0.6 11.3 ± 0.1

TiSiN S3-0 51.4 ± 0.4 1.8 ± 0.1 12.6 ± 0.1 34.2 ± 0.5 –

vlestoi

3

oFttteetsntbfssceoaiotdcw

tVtwpttiba

S3-4 52.9 ± 0.1 1.5 ± 0.1

TiSiVN S3-8 52.8 ± 0.2 2.0 ± 0.1

S3-12 52.4 ± 0.3 2.1 ± 0.2

anadium additions displayed lower hardness and Young’s modu-us values as compared to the other V rich coatings. The hardnessnhance with V additions is probably due to the presence of V inolid solution. In fact, the similar level of residual stresses as func-ion of V content measured, which revealed to be from the rangef 3–4 GPa, and the observed shift of peaks to higher angles with Vncorporation, supports the previous affirmation.

.2. Continuous and isothermal oxidation in air

The effect of increasing Si and V content on the onset point ofxidation of the TiN and TiSiN systems, respectively, are shown inig. 4. Silicon incorporation strongly improved the onset of oxida-ion of the coatings, with the highest silicon content (S3-0) showinghe highest oxidation resistance. Kacsich et al. [34,35] reported thathe high oxidation resistance of these coatings was due to the pres-nce of a protective silicon-rich oxide layer. This layer acted as anfficient diffusion barrier against oxygen and metal ions diffusion;hus, it protected the coating from further oxidation. Our resultshowed that the onset point of oxidation of films decreased sig-ificantly down to a temperature of approximately 500 ◦C (lowerhan TiN) for all coatings containing vanadium. Interestingly, thisehaviour was independent of Si content of the films and, there-ore, the vanadium incorporation interfered with the diffusion ofilicon and the consequent formation of a continuous protectiveilicon oxide layer. Furthermore, all V-containing coatings wereompletely oxidized at a lower temperature than TiN film. Lewist al. [7], Zhou et al. [36], and Franz et al. [37], who studied the effectf V doping on the oxidation behaviour of multilayered TiAlN/VNnd CrAlVN coatings, respectively, also observed a similar decreasen the onset point of oxidation. In order to better characterize thexidation behaviour of the V rich coatings, isothermal oxidationests were conducted at selected temperatures using S2-8 coatingeposited onto FeCrAl alloy substrates. The isothermal curves wereompared with those of TiSiN coating with similar Si content andith TiN (see Fig. 5).

Comparing the mass gain of the coatings and taking into accounthe testing temperature in each case, it is possible to conclude that-free TiSiN coating (S2-0 tested at 900 ◦C) is much more resistant

o oxidation. This coating exhibits a typical parabolic oxidationeight gain as a function of time indicating the formation ofrotective silicon-rich oxide layer, as referred to above. Similarype of evolution is exhibited by TiN coating; however, in this case
he titanium oxide scale does not protect the coatings and, thus,t is effective only at low isothermal temperature [38]. It shoulde remarked that during continuous oxidation test up to 1200 ◦C,t 900 ◦C TiN coatings was already half consumed. Regarding the
10.7 ± 0.0 32.1 ± 0.2 2.7 ± 0.110.7 ± 0.1 27.2 ± 0.2 7.3 ± 0.110.3 ± 0.1 24.9 ± 0.9 10.4 ± 0.1

isothermal curves of TiSiN coating with V additions (S2-8), it canbe concluded that, independently of the isothermal temperatureand time of exposure, their mass gain is always much higher thanTiN and V-free TiSiN coatings (S2-0). The isothermal curves testedat 550 and 600 ◦C showed two steps: at an early stage, the weightgain is rapidly increasing almost linearly up to approximately0.1 mg/cm2 (particularly at 600 ◦C), whereas in the second stage,a parabolic evolution can be fitted. Isothermal annealing of S2-8at 700 ◦C shows significant increase of mass gain following theparabolic evolution.

3.2.1. Surface structureAs it would be expected, both TiN annealed at 600 ◦C and S2-0

coating annealed at 900 ◦C exhibited several intense peaks belong-ing to the (1 1 0), (1 0 1), (2 0 0), (1 1 1) and (2 1 0) of a tetragonalphase related to rutile/TiO2 (ICDD card 76-0649) (not shown here).Weak peaks assigned to TiN (ICDD card 87-0628) were observedin the XRD spectra of the S2-0 coating, which was in a good agree-ment with its annealing curve, since only partial oxidation of thefilm was expected. Moreover, diffraction peaks associated with Sioxide were not detected on the oxidized surface of S2-0 film sug-gesting an amorphous character of the protective oxide layer. Thisresult is in agreement with literature, where similar observationwas reported [34,35].

XRD patterns of S2-8 oxidized coatings in isothermal tests areshown in Fig. 6. The main phases detected by XRD were Ti–V–Oand V–O oxides; however, depending on the isothermal temper-ature and time, different orientations and peaks assigned to V–Owere perceived. After annealing at 550 ◦C during 1 h, the follow-ing phases could be identified: Ti(V)O2 (ICDD card 77-0332) and�-V2O5 (ICDD card 41-1426) with a preferred orientation follow-ing the (0 0 1) plane (peak at 20.3◦), together with the f.c.c. nitridefrom the coating. The same phases were detected at 600 ◦C during10 min, however, �-V2O5 phase showed random orientation. Fur-ther increase in isothermal time (annealing at 600 ◦C during 30 min)resulted in loss of TiN peaks and the orientation of the �-V2O5 phaseis again following (0 0 1) plane. The absence of TiN peaks is in goodagreement with the isothermal curve shown in Fig. 5. Progressiveoxidation is demonstrated in the intensity of (0 0 1) reflection after30 min of annealing, where the amount of �-V2O5 is clearly higher.These changes in phase proportion and orientation have alreadybeen observed by several authors during the oxidation of TiAlN/VNmultilayer and TiAlVN single layered coatings [1,7,8,36]. Increas-
ing the heat treatment to 700 ◦C, which is temperature higher thanthe melting point of �-V2O5 (∼685 ◦C [10,36]), resulted in the ran-dom orientation of this phase; moreover, a sharp peak at 2� = 12.8◦
identified as �-V2O5 (ICDD card 45-1074) with (0 0 2) preferential

118 F. Fernandes et al. / Applied Surface Science 289 (2014) 114– 123

Fig. 2. SEM fracture cross section micrographs of: (a) coating TiN, (b) coating S1-0,(c) coating S1-12.

Fig. 3. (a) Effect of Si content on hardness and Young’s modulus for the TiN system.(b) Effect of V content on hardness and Young’s modulus for coating S1-0.

orientation appearing [39]. �-V2O5 phase is a result of loss of Ofrom �-V2O5, due to a reduction process [40]. We should empha-size here that vanadium oxides observed on the oxidized surface areoptimal from the tribological point of view. Vanadium oxide couldacts as low-friction solid lubricant or, for higher temperatures, asliquid lubricant. Therefore the drop in onset oxidation tempera-ture should not be considered as a significant drawback of TiSiVNcoatings.

3.2.2. Surface morphology and characterizationTypical SEM surface micrographs of the S2-0 and V-rich coating

(S2-8) isothermally oxidized at different temperatures and times

Fig. 4. Thermal gravimetric oxidation rate of coatings deposited on Al2O3 substratesusing the linear-temperature ramp (RT to 1200 ◦C at 20 ◦C/min).

F. Fernandes et al. / Applied Surface

Fig. 5. Thermo gravimetric isothermal analysis of coatings exposed at differenttemperatures.

Ft

awtft

Fa

ig. 6. X-ray diffraction patterns of oxidized surface of: coating S2-8 annealed atemperatures ranging from 550 to 700 ◦C.

re shown in Figs. 7 and 8, respectively. EDS and Raman analysesere carried out on the oxidized surface of the coatings to identify

heir composition. On the oxidized surface of the S2-0 film, two dif-erent zones could be detected, a dark grey phase evenly distributedhroughout the surface and white islands. EDS analysis showed that

ig. 7. SEM observation of surface morphology of coating S2-0, after 1 h oxidationt 900 ◦C.

Science 289 (2014) 114– 123 119

the white oxide was rich in Ti (crystals of TiO2), and the dark greyphase composed mainly by Si and Ti. The Raman spectra of thesezones are plotted in Fig. 9. As can be observed the light grey phasedisplays Raman active modes at 152, 253, 454 and 610 cm−1 [3,23]assigned to rutile, confirming its presence in the oxidized coatingdetected by XRD. The same phase, with much less intensity, wasdetected for the dark grey phase; however, strong additional peaksassigned to anatase (TiO2) were observed. This finding corroboratesthe report of Pilloud et al. [23], who showed that anatase Ramanbands increased with the silicon content in TiSiN films, whereasthat of rutile decreased. As only strong peaks of rutile were detectedby XRD, this indicates that the amount of anatase should be smalland, therefore, not detectable by XRD diffraction. Since Raman pen-etration depth is relatively limited, the absence of silicon oxide inRaman spectra indicated that Si–O was below theTiO2 phase iden-tified by XRD. This finding corroborates the results of Kacsich et al.[34,35]; they observed titanium oxide layer formed at outmost sur-face and silicon oxide sublayer (TiSiN coatings).

The analysis of the oxidized surface of S2-8 coating reveals thatthe surface oxide morphology is different from that of coating S2-0. It displays a floret-like structure formed by light and dark greyzones. At 550 ◦C these phases are tiny distributed throughout thesurface, being difficult to clearly identify the boundaries betweenthem. At a temperature of 600 ◦C, the separation between bothzones is evident with the shape of darker zone suggesting some typeof dendritic growth. With the increase of the isothermal time, anincrease in the area covered by the grey dark phase was observed.The temperature 700 ◦C led to a different oxide morphology, i.e.the appearance of a black zone in some rosettes and the segrega-tion of a white phase to the boundaries of the floret-like structure.This change in the microstructure can be associated to the meltingof �-V2O5, which originates a smoother surface (Fig. 8d)). Fig. 10shows examples of EDS spectra of points identified in Fig. 8 (phases1 and 2). As can be observed, Ti K� peak (4.931 keV) overlaps the VK� peak (4.952 keV) and, therefore, V K� peak should be taken inconsideration for analyzing V importance. In order to identify thedifferent oxide phases marked in Fig. 8, the ratio between the peaksintensities of V K� (Si K�) and Ti K� from EDS analyses are sum-marized in Table 3. To a relative increase in this ratio correspondsthe preferential formation of the oxide of that element. EDS analy-ses of points 1 and 2 reveal similar compositions for the grey lightand dark zones, suggesting the presence of oxides containing Ti, Siand V. However, darker phase has clearly a higher V content (seeTable 3). The Raman spectrum taken at grey dark phase (see Fig. 11)shows the presence of intense peaks assigned to �-V2O5 [3,8,10]and small peaks related to rutile (TiO2). On light grey phase thespectrum is much less defined, with broader bands. �-V2O5 peaksalmost vanished whereas rutile bands are enhanced. Low crystal-lized Ti–O phases give very similar Raman spectra, which evolve,after annealing, either as anatase or rutile [41,42] phases. In sum-mary, the dark grey phase can be assigned mainly to �-V2O5, andsmall quantities of TiO2, probably as a bilayer with �-V2O5 on thetop and TiO2 underneath, as suggested by EDS measurements. Infact, the high ratio of V/Ti of point 3 marked in Fig. 8, allows iden-tifying a V-rich oxide, on the surrounding regions of the dark greyphase. The light grey phase should be attributed to low crystallizedTi(V)O2. These results agree to the XRD results acquired on the oxi-dized surface of the films where both Ti(V)O2 and �-V2O5 phaseswere indexed. When the isothermal time was prolonged from 10 to30 min at 600 ◦C, the increase of the intensity of the �-V2O5 peakobserved in the XRD patterns was is in a good agreement with ahigher amount of grey dark phase observed on the oxidized surface
morphology.
Similarly to S2-0 coating signals from Si–O oxide were neitherdetected by XRD nor by Raman spectroscopy. However, accordingto EDS analysis this oxide should be present. As these signals are

120 F. Fernandes et al. / Applied Surface Science 289 (2014) 114– 123

Fig. 8. Typical surface morphology of oxidized coatings: (a) coating S2-8 exposed to 55exposed to 600 ◦C during 30 min, (d) coating S2-8 exposed to 700 ◦C during 10 min.

Fig. 9. Raman spectra of dark and light grey phases identified in Fig. 7.

TR

able 3atio between the peaks intensities of V K� (Si K�) and Ti K� from EDS analyses of points

1 2

RatiosV K�/Ti K� 0.04 0.1

Si K�/Ti K� 0.77 0.99

0 ◦C during 1 h, (b) coating S2-8 exposed to 600 ◦C during 10 min, (c) coating S2-8

coming from a sub-surface layer, an oxide layer rich in Ti and Sishould coexist with a global chemical composition similar to thatof light grey phase.

From the three different morphological zones of sampleannealed at 700 ◦C for 10 min (Fig. 8d)), EDS analysis showed thatthe two grey phases were mainly rich in vanadium and oxygen. Sig-nals of elements from the substrate were detected in the dark phase,particularly Al, which in conjunction with XRD results, allowedidentifying it as �-V2O5. In fact, it is well known, that the loss of Ofrom �-V2O5, due to a reduction process, leads to its transformationin �-V2O5 [40,43]. The presence of highly oxygen reactive elements,such as aluminium, in the substrate can reduce the molten �-V2O5,subtracting O and promoting the formation of �-V2O5. The Ramanspectra of the dark and light grey phases plotted in Fig. 12 showsimilar Raman patterns with peaks related to V2O5. Both �- and�-V2O5 display similar Raman spectra [44]. The segregated whiteareas located at the boundaries of the grey phase (see ratio Si/Ti ofpoint 5) were rich in Si and Ti, suggesting to be silicon and titanium
oxides.
A good correlation was found between the isothermal oxida-tion curve, XRD measurements and surface morphology of S2-8coating. At temperatures below the melting point of �-V2O5, the

marked in Fig. 8.

Points

3 4 5 6

0.15 0.63 0.01 0.680.19 0.58 0.49 1.00


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igit�igfslt�aosop

ig. 10. Energy dispersive X-ray analysis (EDAX) of points 1 and 2 identified in Fig. 8.

sothermal oxidation curve starts with a linear increase in the massain and then changes to a parabolic law. The XRD spectra of spec-men S2-8 oxidized during 10 min at 600 ◦C, in the first stage ofhe oxidation process, revealed the presence of intense peaks of-V2O5. Further increase in isothermal time to 30 min showed an

ncreased amount of �-V2O5, as confirmed by the increase of therey dark phase, above attributed to V-rich zones. XRD patternor this sample also showed an increase in Ti(V)O2 peaks inten-ity suggesting a thickening of the oxide scale. Therefore, the firstinear part of the isothermal curve is due to the oxidation of bothitanium and vanadium to form Ti(V)O2 and �-V2O5, respectively.-V2O5 is a non-protective oxide and Ti(V)O2 is still not in enoughmounts to form a continuous protective scale, reason why a linear
xidation rate is observed. For longer oxidation times, the progres-ive thickening of the Ti-rich oxide scale will make difficult thexygen diffusion inward retarding the reaction and promoting thearabolic behaviour. However, the protection is not so efficient
Fig. 11. Raman spectra of phase 1 and 2 identified in Fig. 8.

Fig. 12. Raman spectra of phase 4 and 5 identified in Fig. 8.

as in TiN and/or Ti–Si–N samples (see Fig. 5) due to the disrup-tion induced by the outward/lateral diffusion of V ions through theoxide scale. In any condition the protective Si–O layer could be alsoformed. Similar isothermal evolution was reported by Zhou et al.[36], which during isothermal oxidation of multilayered TiAlN/VNcoatings have detected �-V2O5 at the initial stage of oxidation.In their system, the mass gain curves also showed initially a lin-ear evolution, followed by a parabolic growth, similar to whatwas observed here for TiSiVN films. Keller and Douglass [45] alsoshowed that pure vanadium oxidizes linearly, whereas the additionof different elements (Al, Cr, Ti) could transform the initial linearmass gain for a parabolic one due to the modification of the formedoxide layer.

At 700 ◦C, temperature higher than the melting point of �-V2O5, very high oxidation rates were measured and high amountsof V-oxides were detected by XRD. The extremely high parabolicevolution suggests a very high diffusion rate of the reactive ionsthrough the oxide scales.

A comparison between TiSiVN with CrAlVN and TiAlVN (insingle layered or multilayered configuration) coatings, reveals alower onset point of oxidation of approximately 100 ◦C. In CrAlVNand TiAlVN coatings the onset point of oxidation was close to∼600 ◦C [8,36,37]. In the case of the TiAlVN films, signals of lubri-cious �-V2O5 were detected as soon as the oxidation started,while at high temperatures only AlVO4 and TiO2 were identi-fied [8]. For CrAlVN coatings, AlVO4, (Al, Cr, V)2O3 as well as�-V2O5 oxides were observed for an annealing temperature of700 ◦C [37]. In these studies, the lower onset point of oxidationdisplayed by the V-containing coatings was explained by the reac-tions occurring between protective oxides and vanadium (such asformation of Al–V–O phases). In the present TiSiVN system com-bination of V with titanium or silicon oxide was not detected.Therefore, we suggest that the loss of oxidation resistance ofTiSiVN coatings is due to the rapid oxidation of V and the out-wards diffusion of its ions through the Ti-rich oxide scale. Furtherwork is needed to understand the mechanisms involved duringthe oxidation deterioration of TiSiN coatings with V additionsand their corresponding oxide scale formation. The performanceof solid solution coatings for high temperature sliding applica-tion should be investigated in further work. Rapid formation ofvanadium oxide could indeed reduce friction at temperatureshigher than 500 ◦C; nonetheless, the wear resistance of the coat-ing can be significantly compromised by low oxidation resistance.
In future research, we propose to deposit nanocomposite struc-ture of TiSiVN combining nanograins of c-TiVN embedded intoamorphous Si–N matrix, to control the outwards diffusion ofvanadium.

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. Conclusion

This investigation concerned the influence of V additions onhe structure, mechanical properties and oxidation resistance ofi–Si–V–N coatings deposited by DC reactive magnetron sput-ering. These coatings were compared to TiN and TiSiN films.ccording to XRD analyses, all coatings showed an fcc NaCl-typetructure assigned to crystalline TiN. A shift of the peaks to theight was observed with Si and V additions, indicative of a substitu-ional solid solution. Hardness and Young modulus of TiSiN coatingsas insignificantly changed with increasing V content. The onset of

xidation of the coatings decreased with V additions down to tem-eratures as low as 500 ◦C, independently of the Si and V content inhe coatings. TiN and TiSiN coating exhibits a typical parabolic oxi-ation weight gain as a function of time, while a different evolution

s displayed by TiSiVN films. At temperatures below the meltingoint of �-V2O5 (∼685 ◦C) two stages were exhibited: at an earlytage, the weight increase over time is linear, whilst, in a secondtage a parabolic evolution could be fitted to the experimental data;n the other hand, at high temperatures only a parabolic evolutionas fitted. �-V2O5 showed to be the main phase present at the

xidized surface of coatings. Reduction of this phase occurred foremperatures above their melting point. The relative amounts of2O5 detected at the oxidized surface of V rich films are promising

o achieve the envisaged good tribological properties; however itan be significantly compromised by their low oxidation resistance.

cknowledgments

This research is sponsored by FEDER funds throughhe program COMPETE–Programa Operacional Factores deompetitividade–and by national funds through FCT – Fundac ãoara a Ciência e a Tecnologia, under the projects: PEst-/EME/UI0285/2013, CENTRO-07-0224-FEDER-002001 (Maisentro SCT 2011 02 001 4637), PTDC/EME-TME/122116/2010nd Plungetec, as well as the grant (SFRH/BD/68740/2010).

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Annex H


Annex H

F. Fernandes, J. Morgiel, T. Polcar, A. Cavaleiro, Oxidation and diffusion

processes during annealing of TiSi(V)N films, (2014), under review, “Thin Solid

Films”.

Annex H


Oxidation and diffusion processes during annealing of TiSi(V)N films

F. Fernandes1,*

, J. Morgiel2, T. Polcar

3,4, A. Cavaleiro

1

1SEG-CEMUC - Department of Mechanical Engineering, University of Coimbra, Rua Luís Reis

Santos, 3030-788 Coimbra, Portugal.

2Institute of Metallurgy and Materials Science of Polish Academy of Sciences, Krakow, Poland

3National Centre for Advanced Tribology (nCATS), School of Engineering Sciences, University of

Southampton, Highfield, Southampton, SO17 1BJ, UK.

4Department of Control Engineering Czech Technical University in Prague Technicka 2, Prague 6, 166

27 Czech Republic.

*Email address: [email protected], tel. + (351) 239 790 745, fax. + (351) 239 790 701

Abstract

The effect of V additions on oxidation resistance, oxide scale formation and diffusion

processes for TiSiVN system and their comparison to TiSiN is investigated. A dual layer

oxide was formed in the case of TiSiN coating with a protective Si-O layer at an

oxide/coating interface; however, in zones of film defects a complex oxide structure was

developed. V additions increased the oxidation rate of the coatings as a result of the inversion

of the oxidation mechanism due to ions diffusion throughout the oxide scale, which inhibited

the formation of a continuous protective silicon oxide layer.

Keywords: TiSiVN system, structural evolution, oxidation, oxide scale, diffusion processes

1. Introduction

TiSiN solid solution or nanocomposite coatings are well established in applications

like high speed cutting and dry machining processes due to their excellent oxidation resistance

and high hardness. However, the main drawback of the aforementioned coating system is their

high friction coefficient, typically in the range of 0.6 - 1.3 at room and high temperatures [1].

Recently, the formation of thin reaction films in sliding contacts has provided the basis for

development of adaptive, self-lubricating coatings guaranteeing low friction at high

temperature. Magnéli phase oxides based on Ti, Si, Mo, W and V, with low shear strength


Annex H


have attracted the scientific community and particular attention has been given to vanadium

oxide. The beneficial influence of Magnéli oxides formed by oxidation of vanadium on the

friction has already been reported for ternary CrVN [2], (V,Ti)N [3], multilayer AlN/VN [4],

quaternary single layer or multilayered CrAlVN [5-6], and TiAlVN [7-8] systems. Recently

we have reported the effect of V incorporation on the structure, mechanical properties and

oxidation resistance of TiSiVN films deposited by DC reactive magnetron sputtering [9].

Lubricious vanadium oxides have been successfully detected on the oxidized surface of these

films. Here we show the thermal stability, oxide formation and, particularly, diffusion

processes during annealing of TiSiVN films deposited by magnetron sputtering.


TiSiN and TiSiVN coatings with approximately the same silicon content and about 2.5

µm of total thickness, were deposited on FeCrAl alloy substrates by a d.c. reactive magnetron

sputtering machine equipped with two rectangular, Ti (99.9%) and TiSi2 (99.9%), magnetron

cathodes working in unbalanced mode. V incorporation was achieved by inserting 8 pellets of

vanadium into the erosion zone of Ti target. In both cases Ti and TiN adhesion layers were

deposited as bonding layers improving coating to substrate adhesion. The interlayer also

contained V when V pellets are placed in the Ti target. A detail description of the deposition

conditions is reported elsewhere [9]. Temperature effect on the structure of the V rich coating

was characterized by in-situ hot-XRD analysis in the range from 500 ºC up to 750 ºC, in open

air, using a grazing incidence angle of 2º and Co Kα radiation (1.789010 Å). Between each

selected temperature a step of 10 min holding time was allowed for thermal stabilization and

30 min time acquisition was used. Oxidation of films was assessed by thermogravimetric

analysis (TGA) using industrial air (99.99% purity). The coatings were isothermal tested at

different selected temperatures and times, based on the thermogravimetric oxidation curves of

films performed at constant linear-temperature ramp (RT to 1200◦C at 20◦C/min) shown in

reference [9]. These isothermal tests were performed having in consideration the temperatures

where the oxidation showed takes mainly place. After annealing, the cross section thin foils of

oxidized films was prepared by focused ion beam (FIB) and analyzed by transmission

electron microscope (TEM) equipped with an energy-dispersive x-ray (EDS) spectroscopy

system.

Annex H



3.1. Phase evolution during annealing (in situ)

The investigated coatings with chemical composition of Ti0.80Si0.13N and

Ti0.65Si0.11V0.14N showed a typical columnar morphology and fcc NaCl-type structure

assigned to crystalline TiN with Si and V in solid solution [9]. Fig. 1 plots the high

temperature in-situ XRD spectra evolution for TiSiVN coating together with as deposited and

cooled state patterns acquired at room temperature (RT). The first signs of oxides were

detected at 500 ºC and identified as rutile type phase with V in solid solution, the Ti(V)O2

(ICDD card 77-0332). Further increase in temperature to 550 ºC led to an increase of the

Ti(V)O2 peaks intensity and the appearance of a new phase: the α-V2O5 (ICDD card 41-1426)

oxide with orthorhombic symmetry (peak at 23.6º). At 600 ºC a strong increase in intensity of

the α-V2O5 oxide peaks with a strong preferred orientation following the (001) plane was

observed (peak at 23.5º). Moreover, a metastable phase was formed, β-V2O5, as a result of

oxygen loss from α-V2O5 due to a reduction process [10]. It should be also pointed out that

the higher intensity of V2O5 peaks in relation to other peaks suggests higher quantity on the

oxidized volume. The temperature increase to 650 ºC led to TiN peaks vanishing suggesting

the formation of a thick oxide scale. At 675 ºC, the peaks associated with the V2O5 phases

disappeared, which was in a good agreement with the melting point of V2O5 oxide (reported

in the range of 670 – 685 ºC [11]). On the other hand, peaks related to a V3O5 phase, which

resulted from the α-V2O5 reduction, were found. In fact, in case of oxygen deficiency, a shift

to lower valence state leads to the formation of reduced oxides, such as V3O5 identified in the

current XRD pattern. Further increase in temperature resulted mostly in the increase of V3O5

peaks intensities due to continuous reduction of liquid V2O5 phase. Ti(V)O2, V3O5, V2O3 and

VO reduced phases were found after annealing, as well as a new polymorph V2O5 oxide: γ’-

V2O5 (ICDD card 85-2422). This oxide has an orthorhombic structure as α-V2O5 phase, but

with different diffraction planes. These changes could be attributed to distortions on the

structure motivated by the cooling of liquid-phase. Double chains exist in both phases, but

VO5 pyramids alternate up and down individually for γ’- V2O5, whereas they alternate by

pairs for normal α-V2O5 phase [12].

Annex H


Figure 1 - XRD spectra of TiSiVN film at different temperatures (upper) and position of main peaks of

corresponding phases (lower).

3.2. Isothermal oxidation in air, surface morphology and characterization

TiSiN film was annealed at 900 0C for 1 hour, TiSiVN coating was annealled at two

different isothermal temperatures. The results of the thermo gravimetric analysis are presented

in Fig. 2 a). As expected, TiSiN film is much more resistant to oxidation, showing a typical

parabolic oxidation weight gain as a fuction of time, which indicates the presence of

protective oxide scales. Vanadium addition to TiSiN film strongly reduced the oxidation

resistance. The isothermal curves at 550 ºC and 600 ºC showed two steps: they started with a

linear increase in mass gain and then followed with a parabolic evolution. The surface

morphologies of the oxidized coatings are shown in Fig. 2 b-c) for coatings TiSiN tested at

900 ºC and TiSiVN oxidized at 600 ºC, respectively. The detailed description of the surface

oxide constitution, based on XRD diffraction, Raman spectroscopy and SEM-EDS analyses,

can be found in our previous study [9]. Therefore we will only summarize here the main

results to support investigation aimed at diffusion processes and described later. Annealed

TiSiN film (see Fig. 2b) displayed two different surface features: white and dark gray islands

evenly distributed on the surface. Raman analyses revealed that white phase was rutile (TiO2),

while dark gray phase was a mixture of rutile and anatase. However, only rutile peaks were

detected by XRD diffraction suggesting very limited amount of anatase in the dark gray

islands. Furthermore, strong signals of Si were detected on dark zone by EDS. However, the

signals of silicon oxide were neither detected by XRD nor by Raman spectroscopy indicating

16 18 20 22 24 26 28 30 32 34 36 38 40 42 44 46 48 50 52 54

cooling down) (after

VOV

2O

3

V3O

5

'V

V

V

Ti(V)O2

RT

750 0C

700 0C

675 0C

650 0C

600 0C

550 0C

500 0C

2

Rela

tive Inte

nsity

RT

TiN

Annex H


amorphous character, which corroborates previous reports [13-14]. Silicon oxide was

positioned below Ti-O rich layer. Different morphology and oxide phases were detected on

the oxidized surface of TiSiVN coating. At 600 ºC, the film displayed a floret-like structure

formed by light and dark gray phases. XRD diffraction showed the presence of α-V2O5 and

Ti(V)O2 oxides. Although EDS analysis revealed similar spectra for both phases, the much

less intensity of the V peak on the light gray phase, and its conjugation with Raman analysis

allowed identifying dark and light gray phases as α-V2O5 and Ti(V)O2 oxides, respectively.

Similar phases were detected at 550 ºC by XRD; however, only small dark gray areas were

observed at the surface. These results match well with high temperature in-situ XRD

diffractograms (Fig. 1), where these phases were indexed. Similar to TiSiN coating, the signal

from Si-O phase was neither detected by XRD nor by Raman.

Figure 2 - a) Isothermal oxidation curves of coatings, b-c) typical surface morphologies.

0 1200 2400 36000.00

0.05

0.10

0.15

0.20

S2-8 500 ºC 1 h

S2-0 900 ºC 1 h

S2-8 600 ºC 30 min

a)

Weig

ht gain

(m

g/c

m2)

Time (s)

50 µm 500 µm

Dark gray phase

White phase

5 µm

Gray phase

Dark phase

b) c)

20 µm

TiSiN 900 0C 1h TiSiVN 600 0C 30 min

Annex H


Figure 3 shows the growing mechanism of the V2O5 phase over the Ti(V)O2 oxide.

Nucleation points of V2O5 started being formed over the oxidized surface, growing up by

vanadium lateral diffusion as suggested by the V depleted white and V rich black regions

marked in the SEM micrograph. In order to investigate the oxide scale films growth, TEM

cross-sections were prepared by FIB from TiSiN and TiSiVN films annealed at 900 and 600

ºC, respectively.

Figure 3 - Oxidized surface of TiSiVN coating showing the growing mechanism of V2O5 phase.

3.3. Cross section of oxidized surfaces by TEM and oxidation kinetics

Fig. 4 and 5 displays the bright-field STEM images, associated elemental maps and

scanning elemental profiles from cross section of oxidized TiSiN and TiSiVN coatings. In the

case of TiSiN film a multilayer of oxides can be identified, being more complex close to the

film defect (white zones) shown in Fig. 4a): (i) an outer Ti-rich layer comprised by shaped

crystals with bigger size in the top of the film defect, zone corresponding to the white phase

and, (ii) a Si-rich layer, which is itself divided in 3 layers on the zone around the defect, an

intermediate layer containing Ti, sandwiched between two Ti-free layers, being the external

porous and the internal one very compact. Far from the film defect (left zone of the

micrograph), below to the TiO2 crystals only a homogeneous Si-rich layer was observed. The

measured elemental cross-section depth profiles for two lines were plotted in Fig. 4b) and 4c),

respectively. The profiles corroborated STEM/EDX elemental mapping showing in detail the

composition of surface oxides. The surface layer formed exclusively of Ti-O was followed by

a Si-rich layer with scattering in the signals intensity, in agreement with brightness intensity

in figures 4 a). It is clear in Ti-signal a small increase in intensity in the zone of high-Si

content. This variation is more intense in line 2 than in line 1, suggesting an influence of the

film defect, which is closer to line 2. Continuous and compact Si-O layer at the oxide/coating

Floret-like V2O5 phase

100 µm

V depleted white region

Nucleation points

V rich black

region

Annex H


interface acts as an efficient barrier against oxygen and metal ions diffusion and thus

protecting the coating from further oxidation [13-14].

It is evident that Ti oxide forms first on the surface of the film due to the higher

affinity of Ti for O than Si [13-14] and then, with further increase in temperature, silicon

oxide is formed due to the progressive segregation of Si. The oxidation process will then

occur through the inwards diffusion of O2-

through the TiO2 layer and the outwards diffusion

of Ti4+

through the Si-O layer [15]. Nevertheless, due to the presence of film defects, in the

first stage of the oxidation process, a high amount of Ti ions will be supplied in that zone

corresponding to the oxidation occurring in the defect walls. This process will promote the

formation of a porous Si-O layer, with low diffusion barrier performance to the out diffusion

of Ti4+

ions, leading to large TiO2 crystals on the oxide scale surface. From the moment that a

Si-O barrier layer is formed in the defect walls (see Si signal in figure 4 a)) the oxidation

process will be controlled by the O2-

inward and Ti4+

outward ions diffusion trough the Si-O

layer [15]. It should be remarked that this phenomenon should not influence significantly the

global oxidation behavior of the Ti-Si-N coating (which shows the parabolic behavior

presented in Fig. 2 a), in agreement with literature [13-15]) since, on the one hand, it only

occurs in a few defects in the films surface (see Fig. 2 b)) and, on the other, it should be more

intense in the first stage of the oxidation, during heating up to the 900 ºC isothermal oxidation

temperature.

Figure 4 - a) Bright field STEM/EDX maps of TiSiN coating oxidized at 900 0C during 1h. b) and c). Elemental

profiles along the cross section of the oxidized coating from: b) line 1, c) line 2.

O Si TiSTEM

Gray phase

White phase

a)

Lin

e 1

Lin

e 2

2 µm

Annex H



TiSiVN coating exhibited, after annealing, two-layers oxide structure with a thick

porous inner layer and an outer discontinuous layer of well-defined crystals; the latter

corresponded to the gray areas described above. The elemental maps shown in Fig. 5a)

suggested that the majority of vanadium was located in the outer crystal layer, whereas inner

layer was Ti and Si rich. This result corroborates the detection of the rutile type compound

Ti(V)O2 indexed by XRD and Raman spectroscopy in the light gray phase in Fig. 2 [9].

Elemental lines shown in Fig. 5b) and 5c) showed that diffusion of V occurred exclusively

within the oxidized volume, since a constant signal was measured for the remaining non-

oxidized coating. In general, we observed a thicker non oxidized layer in the regions covered

by V-O crystal phase. Furthermore, comparing either the V content integrated intensities of

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.50.0

0.4

0.8

100

200

300

400

Ti K

O K / Ti K

Si K / Ti K

Su

bstr

ate

Ti in

terla

ye

r

TiN

in

terla

ye

r

Coating

b)

Oxi

dize

d zo

ne

Counts

Distance (m)

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.50.0

0.4

0.8

50

100

150

200

Ti K

O K / Ti K

Si K / Ti K

Su

bstr

ate

Ti in

terla

ye

r

TiN

in

terla

ye

r

Coating

c)

Oxid

ized z

one

Counts

Distance (m)

Annex H


Fig. 5a) and b) or the brightness of V signal between the right and left part of Fig. 5 a), it can

be concluded that V has to be diffused from the left to the right zone of this figure. Finally, it

should be remarked that there is no indication of formation of any Si-rich layer, being Si

signal uniformly distributed in both oxide scale and non-oxidized coating. Therefore, a dense

compact protective silicon oxide layer localized in subsurface was not formed being Si-O

randomly distributed in the Ti(V)O2 porous scale.

Figure 5 - a) Bright field STEM/EDX maps of TiSiVN coating oxidized at 600 0C during 30 min. Elemental

profiles along the cross section of the oxidized coating from: b) line 1, c) line 2.

Ti Si VSTEM

Light gray phase

Dark gray phase

Lin

e 2

Lin

e 1

2 µma)

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.50.0

0.4

0.8

400

800

1200

1600

2000

Ti K

O K / Ti K

Si K / Ti K

V K / Ti K

Su

bstr

ate

Ti in

terla

ye

r

TiN

in

terla

ye

r

Coating

b)

Oxidized zone

Counts

Distance (m)

Annex H



Isothermal curve of TiSiVN coating oxidized at temperatures below the melting point

of V2O5 started with a linear increase in mass gain. This step, where a continuous Ti(V)O2

oxide layer was grown (see Fig. 1), should be attributed to the conjugation of a fast rate of

mass transport in the oxide with a slow surface step accompanying the gas dissociation,

adsorption and subsequent entry of oxygen into the Ti(V)O2 lattice. In fact, at the very

beginning of the oxidation process, due to the high Ti content, a TiO2 layer starts growing.

The presence of V ions and its high solubility with Ti promote the formation of Ti(V)O2 solid

solution, which comprises vanadium cations with lower oxidation states as V3+

ions [16-17].

The substitution of Ti4+

by V3+

ions in the TiO2 lattice would increase the concentration of

interstitial metallic Ti4+

ions and decrease the number of excess electrons. The oxidation

mechanism is inverted in relation to Ti-N, being the outwards diffusion of metallic ions (Ti3+

Tin+

) more favorable that the inwards diffusion of O2-

ions. There are several consequences of

this inversion: (i) the oxidation rate increases substantially, (ii) the kinetics of the oxidation is

not controlled any more by a diffusion process but by the rate of dissociation, adsorption and

combination at the surface of O2-

with metallic ions, and (iii) continuous and compact Si-rich

layer cannot be formed. In fact, the growing of Ti(V)-O on the external part of the scale

leaves a less compact zone in the interface between the non-oxidized material and the oxide

scale, facilitating the local segregation of Si inside the Ti(V)-O instead its segregation for the

interface. The elemental maps suggest that Si-O and Ti(V)O2 coexist in the inner porous

layer, being unable to protect the material from oxidation. As a consequence, the oxidation

process is accelerated, which is demonstrated by the initial rapid oxidation observed in the

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.50.0

0.4

0.8

400

800

1200

1600

2000

Increasing of V

and O counts

Ti K

O K / Ti K

Si K / Ti K

V K / Ti K

Su

bstr

ate

Ti in

terla

ye

r

TiN

in

terla

ye

r

Coating

c)

Oxidized zone

Counts

Distance (m)

Annex H


isothermal curve. Thongtem et al. [16] studied the effect of V doping to Ti-based alloys and

observed similar behavior with increasing V content.

The experimental evidence shows that V2O5 crystals were formed exclusively on the

surface. When V3+

or V4+

arrive to the surface, they are further oxidized to V5+

combining

with O ions and creating nucleation points for the V2O5 growth (see Fig. 3), which will

expand and grow up by vanadium lateral diffusion. The process develops leading to the

formation of the observed floret-like dispersive V2O5 phases that can comprise both α-V2O5

and β-V2O5 phases as evidenced in the XRD patterns at high temperature. As soon as V2O5

oxide expands laterally, the surface will be progressively covered by this oxide making more

and more difficult the access of both types of ions, up to the moment that the rate of

dissociation, adsorption and combination of O2-

with metallic ions is not anymore the kinetics

controlling step but the ion diffusion through the oxide scale. Therefore, after this moment the

isothermal curve will follow a parabolic law (see Fig. 2a)). It should be remarked that V2O5 is

more protective than (Ti,V)O2 since the non-oxidized coating underneath it is thicker.

4. Conclusion

In this investigation, the effect of V additions on oxidation resistance, oxide scale

formation and diffusion processes for TiSiVN system and their comparison to TiSiN was

investigated. In summary, we demonstrated that oxidation of TiSiN is controlled by the

formation of a dense and compact silicon oxide acting as a diffusion barrier, which can be

affected by defects in the as-deposited film. The presence of V3+

ions in the TiO2 scale, during

the oxidation of TiSiVN film, was responsible for an inversion in the ions diffusion process

through the oxide scale leading to, simultaneously, a significant increase of the oxidation rate

and the inability of formation of a Si-O protective oxide.

Acknowledgments



para a Ciência e a Tecnologia –, under the project: PTDC/EME-TME/122116/2010, as well as

the grant (SFRH/BD/68740/2010). T. Polcar acknowledges support from the project MSM

6840770038.

Annex H


References

[1] Y.H. Cheng, T. Browne, B. Heckerman, Mechanical and tribological properties of

nanocomposite TiSiN coatings, E.I. Meletis, Surf. Coat. Technol., 204 (2010) 2123-2129.

[2] M. Uchida, N. Nihira, A. Mitsuo, K. Toyoda, K. Kubota, T. Aizawa, Friction and wear

properties of CrAlN and CrVN films deposited by cathodic arc ion plating method, Surf.

Coat. Technol., 177–178 (2004) 627-630.

[3] J.H. Ouyang, S. Sasaki, Tribo-oxidation of cathodic arc ion-plated (V,Ti)N coatings

sliding against a steel ball under both unlubricated and boundary-lubricated conditions, Surf.

Coat. Technol., 187 (2004) 343-357.

[4] J.-K. Park, Y.-J. Baik, Increase of hardness and oxidation resistance of VN coating by

nanoscale multilayered structurization with AlN, Mater. Lett., 62 (2008) 2528-2530.


Mitterer, Influence of phase transition on the tribological performance of arc-evaporated

AlCrVN hard coatings, Surf. Coat. Technol., 203 (2009) 1101-1105.

[6] Y. Qiu, S. Zhang, J.-W. Lee, B. Li, Y. Wang, D. Zhao, Self-lubricating CrAlN/VN

multilayer coatings at room temperature, Applied Surface Science, 279 (2013) 189-196.

[7] Q. Luo, Temperature dependent friction and wear of magnetron sputtered coating

TiAlN/VN, Wear, 271 (2011) 2058-2066.

[8] K. Kutschej, P.H. Mayrhofer, M. Kathrein, P. Polcik, C. Mitterer, Influence of oxide

phase formation on the tribological behavior of Ti–Al–V–N coatings, Surf. Coat. Technol.,

200 (2005) 1731-1737.

[9] F. Fernandes, A. Loureiro, T. Polcar, A. Cavaleiro, The effect of increasing V content on

the structure, mechanical properties and oxidation resistance of Ti-Si-V-N films deposited by

DC reactive magnetron sputtering, Applied Surface Science, 289 (2014) 114-123

Annex H


[10] A. Bouzidi, N. Benramdane, S. Bresson, C. Mathieu, R. Desfeux, M.E. Marssi, X-ray

and Raman study of spray pyrolysed vanadium oxide thin films, Vibrational Spectroscopy, 57

(2011) 182-186.

[11] R. Franz, C. Mitterer, Vanadium containing self-adaptive low-friction hard coatings for

high-temperature applications: A review, Surf. Coat. Technol., 228 (2013) 1-13.

[12] J.M. Cocciantelli, P. Gravereau, J.P. Doumerc, M. Pouchard, P. Hagenmuller, On the

preparation and characterization of a new polymorph of V2O5, Journal of Solid State

Chemistry, 93 (1991) 497-502.

[13] T. Kacsich, S. Gasser, Y. Tsuji, A. Dommann, M.A. Nicolet, A. Nicolet, Wet oxidation

of Ti34Si23B43, Journal of Applied Physics, 85 (1999) 1871-1875.

[14] T. Kacsich, M.A. Nicolet, Moving species in Ti34Si23N43 oxidation, Thin Solid Films,

349 (1999) 1-3.

[15] F. Deschaux-Beaume, N. Frety, T. Cutard, C. Colin, Oxidation modeling of a Si3N4-TiN

composite: Comparison between experiment and kinetic models, Ceram. Int., 35 (2009) 1709-

1718.

[16] S. Thongtem, T. Thongtem, M. McNallan, High-temperature nitridation and oxidation of

Ti-based alloys, Surface and Interface Analysis, 32 (2001) 306-309.

[17] A. Glaser, S. Surnev, F.P. Netzer, N. Fateh, G.A. Fontalvo, C. Mitterer, Oxidation of

vanadium nitride and titanium nitride coatings, Surface Science, 601 (2007) 1153-1159

Annex I


Annex I

F. Fernandes, T. Polcar, A. Cavaleiro, Tribological properties of self-lubricating

TiSiVN coatings at room temperature, (2014), accept for publication, “Surface and

Coatings Technology”, DOI: 10.1016/j.surfcoat.2014.10.016



Annex I


Tribological properties of self-lubricating TiSiVN coatings at room temperature

F. Fernandes1*

,T. Polcar2,3

, A. Cavaleiro1

1SEG-CEMUC - Department of Mechanical Engineering, University of Coimbra, Rua Luís

Reis Santos, 3030-788 Coimbra, Portugal.

2Department of Control Engineering Czech Technical University in Prague Technicka 2,

Prague 6, 166 27 Czech Republic.

3n-CATS University of Southampton Highfield Campus SO17 1BJ Southampton, UK.

*Email address: [email protected], tel. + (351) 239 790 745, fax. + (351) 239 790

701

Abstract

In the last years, vanadium rich coatings have been introduced as possible candidates

for self-lubrication due to their optimum tribological properties. In the present investigation,

the influence of V incorporation on the wear performance of TiSiN films deposited onto WC

substrates by d.c. reactive magnetron sputtering is reported. The results achieved for TiSiVN

films were compared and discussed in relation to Ti0.80Si0.15N, TiN and Ti0.82V0.15N coatings

prepared as references. The tribological properties of the coatings were evaluated at room

temperature on a pin on disc tribometer equipment using two different counterparts: Al2O3

and HSS balls. The wear tracks, ball-wear scars and wear debris were characterized by

scanning electron microscopy with energy dispersive X-ray spectroscopy (SEM-EDS).

Tribological tests indicated that the wear rate and the friction coefficient of Ti0.80Si0.15N

coating decreased with continuous increase of V content being the overall behaviour strongly

dependent on the counterpart ball material. For Al2O3 balls the wear rate and friction

coefficient of coatings were much lower compared to sliding against HSS steel balls.

Ti0.80Si0.15N showed the lowest wear resistance among all tested coatings, independently of

the counter-body. For V rich coatings tested with Al2O3 balls the polishing wear mechanism

was observed, whereas adhesion wear took place when tested against HSS balls.

Keywords: TiSiVN coatings, Tribology, Wear resistance, Lubricious oxides


Annex I


1. Introduction

Over the past years, ternary titanium-based nitride coatings, such as TiXN (X = B, Cr,

Al, Si, Cr, C, etc.), have been largely employed in industrial applications involving cutting

tools, wear protection, machinery components and high temperature parts, due to their high

hardness and suitable wear and oxidation resistance [1, 2]. However, in dry machining

operations, the tools experience extremely harsh conditions combining abrasion, adhesion,

oxidation, mechanical and thermal loads, which limit their lifetime. Since low friction

coefficients can effectively reduce the contact temperatures during sliding and, consequently,

improve the tribological behavior of tools, a lot of efforts has been carried out in the last years

to develop self-lubricant coatings. These films should retain beneficial properties of (TiX)N

films and offer lubricity through the formation of low-friction surface oxides. Recently, the

attention of the scientific community has been focused on the formation of thin reaction films

of the so-called Magnéli oxide phases based on Ti, Si, Mo, W and V, which possess easy

shear planes [3, 4]. Among these elements, particular attention has been given to vanadium-

containing coatings which form VnO3n-1 Magnéli phases. Over the whole ternary systems the

effect of V doping was only extensively studied for ternary CrAlN [5, 6] and TiAlN [2, 7-9]

coatings in single layer or multilayered configurations. In all the cases a beneficial influence

of V additions on the tribological behaviour of the coatings at room and high temperatures has

been demonstrated. Since other ternary coating systems with superior mechanical properties

have been successfully used as protective coatings in tribological applications, the effect of

vanadium doping on their structure should also be considered. This is the case of TiSiN

system, which shows similar levels of oxidation resistance as CrAlN and TiAlN films.

Preliminary studies on the effect of V additions on the structure, mechanical properties and

oxidation resistance of TiSiN systems have been carried out in our previous investigation

[10]. Hence, as a way to complement that study, the aim of this work was to characterize the

tribological behaviour at room temperature of TiSiVN sputtered coatings, sliding against

Al2O3 and HSS balls. The influence of V addition on the TiSiN system was compared with

the tribological results of TiN, Ti0.82V0.15N and Ti0.80Si0.15N coatings prepared and tested as a

reference.

Annex I


2. Experimental Procedure

TiSiVN coatings with dissimilar Si and V contents, evaluated in our previous work

[10], were deposited by d.c. reactive magnetron sputtering from two rectangular (100×200

mm) magnetron cathodes working in unbalanced mode. High speed steel (AISI M2), FeCrAl

alloy, cemented carbides, silicon and stainless steel substrates were coated for mechanical,

structural, chemical composition and residual stress measurements. A high purity Ti (99.9%)

target, with 10 mm diameter holes, and a high purity TiSi2 (99.9%) target were used in the

depositions. The coatings comprise three TiSiVN coatings with different V and Si contents

and three other coatings deposited as a reference: TiN, Ti0.82V0.15N and Ti0.80Si0.15N. V

incorporation was achieved by filling Ti holes with 4 and 8 V pellets being the remaining

empty holes filled with pellets of Ti. Silicon content was varied by changing the power

density applied to each target. Ti and TiN adhesion layers of approximately 0.24 and 0.45 µm,

respectively, were deposited below final TiSi(V)N coatings. The interlayers also contained V

when V pellets were placed in the Ti target. A negative bias of 50 V was applied to the

substrate holder. All the depositions were performed at a total working gas pressure of 0.3 Pa.

The deposition time was set in order to produce coatings with approximately 2.5 µm of total

thickness, including interlayer referred to above. A full description of the deposition

procedure is given elsewhere [10]. Summary of the deposition parameters used for the

coatings production and their chemical composition determined by electron probe

microanalysis (EPMA) are shown in Table 1 and Table 2, respectively. The denomination of

the samples, as presented in tables 1 and 2, will be adopted throughout the paper in order to

help the coatings identification and the paper reading. The subscript in Ti, Si and V letters

represents the ratio between the at.% of each element in the coating in relation to the at.% of

Nitrogen.

The structure of the films was analyzed by X-ray diffraction (X’ Pert Pro MPD

diffractometer) using a grazing incidence angle of 1º with Cu Kα1 radiation (λ = 1.54060 Å).

The adhesion of the films on the M2 steel substrate was evaluated by a scratch-test apparatus.

Specimens were scratched as the normal force was increased linearly from 5 to 70 N, using a

Rockwell C indenter with a spherical tip with a radius of 0.2 mm, a scratch speed of 10

mm/min and a loading speed of 100 N/min. Critical loads were determined by optical

microscope analysis according to the standard for scratch-test evaluation [11]. For each

specimen, the indicated critical loads results from the average of 3 different scratches. The

hardness and Young’s modulus of films were evaluated by depth-sensing indentation

Annex I


technique (Micro Materials NanoTest) using a Berkovich diamond pyramid indenter. In order

to avoid the effect of the substrate, the applied load (10 mN) was selected to keep the

indentation depth less than 10% of the coating´s thickness. The residual stress values were

calculated through the Stoney equation [12] by measuring the differential deflection between

an uncoated and coated stainless steel 25 mm diameter disk.

The tribological behaviour of the coatings was evaluated by dry-sliding tests at room

temperature on a pin on disc tribometer. Al2O3 and HSS steel balls (both with a diameter of 6

mm) were used as counterparts. The radius of the wear tracks was set to 5.3 and 4 mm for

Al2O3 and HSS steel balls, respectively. All the measurements were performed with a linear

speed of 0.10 m.s-1

, load of 5 N, relative humidity 48±5% and 5000 cycles. The wear rate of

the coatings was determined from the area of the wear track cross section using a 3D optical

profilometer; an average from four different locations was used. The ball wear rates were

evaluated from the wear cap images taken in an optical microscope. To ensure the

reproducibility of results a set of three tests was performed under identical conditions for each

coating. After the wear tests, the wear tracks and wear debris were characterized by scanning

electron microscopy with energy dispersive X-ray spectrometry (SEM-EDS).

Table 1 - Deposition parameters of the coatings.

Target power density

(W/cm2)

Gas flow

(sccm) Deposition time

(min)

Pellets of V at the Ti

target Sample Ti TiSi2 Ar N2

Ti interlayer 10 - 35 - 5 -

TiN interlayer 10 - 35 17 15 -

TiN 10 - 35 17 73 -

Ti0.82V0.15N 10 - 35 17 73 8

Ti0.80Si0.15N 6 1.5 30 17 193 0

Ti0.75Si0.11V0.03N 6 1.5 30 17 193 4

Ti0.65Si0.11V0.15N 6 1.5 30 17 193 8

Ti0.52Si0.20V0.14N 5 2.5 30 17 153 8

Annex I


Table 2 - Chemical composition of films in at.%.

Sample Chemical composition (at.%)

Ti Si V O N

TiN 49.6 ± 0.4 - - 0.38 ± 0.12 49.79 ± 0.13

Ti0.82V0.15N 41.3 ± 0.2 - 7.81 ± 0.11 0.53 ± 0.12 50.42 ± 0.10

Ti0.80Si0.15N 41.3 ± 0.3 6.68 ± 0.03 - 0.55 ± 0.06 51.5 ± 0.2

Ti0.75Si0.11V0.03N 39.3 ± 0.4 5.71 ± 0.06 1.6 ± 0.2 1.22 ± 0.10 52.1 ± 0.2

Ti0.65Si0.11V0.15N 33.6 ± 0.3 5.64 ± 0.05 7.6 ± 0.2 1.42 ± 0.13 51.71 ± 0.14

Ti0.52Si0.20V0.14N 27.2 ± 0.2 10.73 ± 0.06 7.32 ± 0.05 1.96 ± 0.11 52.8 ± 0.2


3.1. Structure and mechanical properties of films

The X-ray diffraction patterns of the as-deposited films are shown in Figure 1. All the

coatings presents an fcc NaCl-type structure assigned to crystalline TiN phase, but different

peaks position and intensities can be perceived. All films presented well defined peaks except

for the coating with higher silicon content (Ti0.52Si0.20V0.14N) for which a strong decrease of

the peaks intensity as well a significant broadening was observed. This result suggests a

decrease of the crystallinity and grain size, in good accordance with literature where for TiSiN

films with high Si contents amorphization of the structure is usually observed [13]. A closer

analysis of the peaks positions indicated that in all cases a fcc solid solution should be

expected. In fact, starting with the individual addition of V and Si to TiN, the peaks are

shifted to higher angles, behaviour associated to a smaller unit cell (insert in Figure 1). Taking

into account that the residual stresses are very similar among these coatings (see table 3), such

a result can only be explained by the presence of V or Si in substitutional solid solution, due

to their smaller radius (0.143 and 0.117 nm, respectively) in comparison with Ti (0.147 nm).

When both Si and V are added to TiN, a synergetic effect of both elements are observed being

the diffraction peaks shifted to further higher diffraction angles. A detailed description of the

effect of Si and V additions on the structure of TiN is shown in our previous work [10].

Annex I


Figure 1 - Typical X-ray diffraction patterns of the dissimilar coatings.

Table 3 shows the hardness, Young´s Modulus, residual stress and scratch test critical

load values of the coatings. V incorporation at the TiN and TiSiN systems slightly increased

the hardness and Young’s modulus value of coatings. Ti0.52Si0.20V0.14N coating exhibited the

lowest hardness and Young´s modulus as compared to the other films, a consequence of the

above referred loss of crystallinity. Residual stresses were similar for all coatings

(compressive; 3 – 4 GPa) suggesting that the shift of TiN peaks to higher angles should be

exclusively attributed to the presence of V and/or Si in solid solution. Moreover, by similar

reasons, the small enhancement in the hardness achieved with V additions is probably due to

solid solution hardening. Similar correlations could be drawn for TiSiN system in comparison

to TiN coating. In relation to the adhesion of films only two failure modes were observed on

the scratch tracks of the coatings: the first cracking (Lc1) and the first chipping (Lc2). Globally,

all the coatings are well adherent to the substrate exhibiting high Lc values with Ti0.82V0.15N

and Ti0.52Si0.20V0.14N being better and worse performing, respectively. In general, Si

incorporation in TiN system globally decreased the critical load values of the coatings.

30 40 50 60 70 80

36 37 38

TiN (111)

TiN (311)TiN (220)TiN (200)

TiN (111)

Rela

tive Inte

nsity

2

Ti0.52

Si0.20

V0.14

N

Ti0.65

Si0.11

V0.15

N

Ti0.75

Si0.11

V0.03

N

Ti0.80

Si0.15

N

Ti0.82

V0.15

N

TiN

TiN

Ti0.82

V0.15

N

Ti0.80

Si0.15

N

Ti0.75

Si0.11

V0.03

N

Ti0.65

Si0.11

V0.15

N

Ti0.52

Si0.20

V0.14

N

Re

lative

In

ten

sity

2

Annex I


Table 3 –Mechanical properties and residual stresses of the dissimilar coatings.

Specimen

Nano-

hardness

(GPa)

Young's Modulus

(Gpa)

Residual stresses

(GPa)

Adhesion critical loads (N)

First coating

cracking Lc1

First chiping

Lc2

TiN 23.2 ± 1.8 324 ± 16 3.3 58.2± 1.4 > 70

Ti0.82V0.15N 24 ± 2 398 ± 21 3.7 62.1± 1.2 > 70

Ti0.80Si0.15N 27± 2 307 ± 10 3.4 46.0 ± 1.4 62.6± 1.1

Ti0.75Si0.11V0.03N 27 ± 3 327 ± 15 3.7 55.2 ± 1.1 64.6 ± 1.1

Ti0.65Si0.11V0.15N 28 ± 2 328 ± 7 4.1 53.6 ± 1.3 64.8 ± 1.3

Ti0.52Si0.20V0.14N 26± 2 288 ± 12 3.2 42.6 ± 1.1 60.6 ± 1.2

3.2 Tribological behaviour

3.2.1 The wear rate of coatings and balls

The wear rate of coatings tested at room temperature were dependent on the

counterpart material, as shown in Figure 2. Generally, when Al2O3 balls were used as

counterparts, the wear rate of coatings was much lower compared to sliding against HSS steel

balls. For reference coatings tested with Al2O3 balls, the tribological behaviour of TiN coating

was improved when vanadium was added; on the other hand, the addition of Si had a negative

effect. As it will be shown later, there is a correlation between the wear rates and the cracking

critical load values of these coatings, which suggests a possible relationship with the fracture

toughness of the coatings and its influence on the wear debris production.

When Si and V were added simultaneously, V had predominant influence. In fact, any

of the TiSiVN coatings showed lower wear rate than the TiN reference, with similar values to

Ti0.82V0.15N sample. It should be remarked that these coatings have, generally, similar Lc

values as Ti0.82V0.15N, i.e. higher Lc than the ones measured for Ti0.80Si0.15N . A similar trend

has been observed for other nitride coating systems containing V [7, 14]. Besides the

influence of V on the Lc values, the high wear resistance of V-rich coatings can also be

attributed to the formation of V lubricious oxides on the wear track of coatings (see further

discussion below). This can explain why Ti0.52Si0.20V0.14N shows lower wear rate than TiN

(and Ti0.80Si0.15N) in spite of its lower Lc value.

The wear rate of the coatings increased significantly when tested against HSS balls,

particularly for the coatings where the V-rich lubricious oxide was identified during the

contact with Al2O3 balls. The lower hardness of HSS ball associated with an easier

incrustation of hard wear debris in the sliding contact promoted abrasion of the wear track

Annex I


with increasing wear rates. Limited formation of vanadium oxides led to higher friction

coefficient, more energy dissipated and more oxide particles in the wear track, as it will be

discussed below.

As expected, the wear rate of the Al2O3 balls is much lower than that of softer HSS

balls. Different trends were found between the wear rates of the coatings and the balls: (i)

HSS ball wear rate increased when the coating wear decreased, and (ii) the wear rate of the

Al2O3 balls was very low and almost constant for all the tests. In both cases, the ball surfaces

were covered with the wear debris; however, HSS balls show large areas with shallow

scratches parallel to the movement of the ball and higher amount of adhered material.

Figure 2 - a) Wear rates of coatings tested against Al2O3 and HSS balls. b) Wear rates of Al2O3 and HSS

counterparts.

0

2

4

6

8

10

12

14 Wear track

Al2O

3 ball

HSS ball

a)

Wear

rate

(10

-6 m

m3/N

.m)

TiN Ti

0.82V

0.15N Ti

0.80Si

0.15N

Ti0.75

Si0.11

V0.03

N

Ti0.65

Si0.11

V0.15

N

Ti0.52

Si0.20

V0.14

N

0.0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8 Wear track

Al2O

3 ball

HSS ball

b)

Wear

rate

(10

-6 m

m3/N

.m)

TiN Ti

0.82V

0.15N Ti

0.80Si

0.15N

Ti0.75

Si0.11

V0.03

N

Ti0.65

Si0.11

V0.15

N

Ti0.52

Si0.20

V0.14

N

Annex I


3.2.2 Friction coefficient

Figure 3 a) and b) shows the friction coefficient of the coatings tested against Al2O3

and HSS balls, respectively. As the radius of the wear tracks used for Al2O3 and HSS balls

tests were not the same and consequently, the sliding distances were different, the evolution

of the friction coefficient is plotted as a function of the total number of cycles (5000 cycles).

All friction curves exhibited two stages: running-in and steady state. Generally, the coatings

showed lower friction coefficient values against Al2O3 showing thus similar trend to the wear

rate. The highest friction coefficient was obtained for Ti0.80Si0.15N coating with values of 1.09

and 1.15 when tested against Al2O3 and HSS balls, respectively; it should be pointed out that

the wear rate was the highest as well. TiN coating presented also a high friction coefficient

when tested against Al2O3 ball, with a large variation in the values with the number of cycles.

Addition of Si and V had a significant impact on the frictional properties when Al2O3

ball was used as a counterpart. Compared to Ti0.80Si0.15N, the friction dropped to one half

when sufficient amount of vanadium was added and silicon content was kept low; the friction

of Ti0.65Si0.11V0.15N is below 0.5 being the lowest together with Ti0.82V0.15N coating. Testing

with HSS balls showed only limited influence of coating composition on the friction

coefficient. Again TiSiVN coatings exhibited the lowest friction, although it was only about

10% lower than that of Ti0.80Si0.15N and similar to that of TiN.

Figure 3 - Friction coefficient vs number of cycles of coatings tested against: a) Al2O3 balls, b) HSS steel balls.

0 1000 2000 3000 4000 50000.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

TiN Ti0.75

Si0.11

V0.03

N

Ti0.82

V0.15

N Ti0.65

Si0.11

V0.15

N

Ti0.80

Si0.15

N Ti0.52

Si0.20

V0.14

N

a)

Friction c

oeffic

ient

Number of cycles

Annex I



3.2.3 Wear mechanisms: Sliding against Al2O3 balls

The changes in the friction coefficient and wear rates can be only partially correlated

with the structural and mechanical properties of the films. To identify the dominant wear

mechanisms scanning electron microscopy with energy dispersive X-ray spectroscopy (SEM-

EDS) was used to analyze the worn surfaces and the wear debris produced in the contact.

Typical 2D profiles of the wear tracks of the coatings tested against Al2O3 are depicted in

Figure 4 a). Profilometer measurements showed that with the addition of V to TiN and

Ti0.80Si0.15N coatings the wear track became shallower. Ti0.80Si0.15N and TiN showed the

deepest and widest wear tracks among all the coatings, which gave rise to their highest wear

rates shown in Figure 2. The wear track of TiN coating was covered by fine scratches parallel

to the relative sliding movement suggesting abrasion as the dominant wear mechanism. The

scratches ended by the accumulation of the dragged material, which can well explain the

fluctuations observed in the friction curve (Figure 5)). EDS analysis performed in the wear

track revealed a mixture of as-deposited and oxidized coating material. Ti0.80Si0.15N coating

showed a similar wear mechanism as TiN; however, due to lower fracture toughness the worn

particles were larger and scratches deeper. The presence of large non-oxidized particles

originated from the coating led to an intensive abrasion of TiN and, particularly, Ti0.80Si0.15N

coating.

0 1000 2000 3000 4000 50000.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

TiN Ti0.75

Si0.11

V0.03

N

Ti0.82

V0.15

N Ti0.65

Si0.11

V0.15

N

Ti0.80

Si0.15

N Ti0.52

Si0.20

V0.14

N

b)

Friction c

oeffic

ient

Number of cycles

Annex I


Figure 4 - 2D profiles of the wear track of coatings tested against: a) Al2O3, b) HSS.

Figure 5 - Wear track of TiN film tested against Al2O3 ball.

Vanadium-containing coatings exhibited a different wear mechanism. Polishing wear

occurred in Ti0.82V0.15N, Ti0.75Si0.11V0.03N and Ti0.65Si0.11V0.15N coatings as demonstrated by

smooth worn surfaces (Figure 6 a)) whereas stick-slip adhesion wear occurred in

Ti0.52Si0.20V0.14N coating resulting in rough surface (Figure 6 b)). The wear track of

Ti0.65Si0.11V0.15N coating (representative of polishing wear), see Fig. 6 c), showed small black

wear debris with vermicular morphology distributed throughout the wear surface. EDS

analysis performed at these well adhered wear debris particles showed that they are mainly

composed of V and O with traces of Ti and Si, suggesting the presence of V-O, Ti-O and Si-O

-0.4 -0.2 0.0 0.2 0.4

0

1

2

3

4

5

6

7

8

-0.4 -0.2 0.0 0.2 0.4

0

1

2

3

4

5

6

7

8D

ista

nce

(

m)

Dis

tan

ce (

m)

b)a)

Tested against HSS

Distance (mm)

Tested against Al2O

3

Ti0.52

Si0.20

V0.14

N

Ti0.65

Si0.11

V0.15

N

Ti0.75

Si0.11

V0.03

N

Ti0.80

Si0.15

N

Ti0.82

V0.15

N

TiN

Ti0.52

Si0.20

V0.14

N

Ti0.65

Si0.11

V0.15

N

Ti0.75

Si0.11

V0.03

N

Ti0.80

Si0.15

N

Ti0.82

V0.15

N

TiN

Distance (mm)

Material accumulation

Scratches

20 µm200 µm

Annex I


oxide phases. It should be remarked that all the other zones of the wear track showed EDS

spectra identical to the as deposited films. After the test with Ti0.52Si0.20V0.14N coating, the

wear scar of the Al2O3 ball was covered with wear debris forming a transfer layer identified

by EDS analysis as a mixture of V-O, Ti-O and Si-O oxides. Friction induced oxidation of

wearing surfaces containing vanadium leading to formation of V-O, which acts as a solid

lubricant in sliding contact [2, 7, 8]. In this study, thick tribolayer consisting mostly of

vanadium oxide was indeed formed on the surface of Ti0.82V0.15N coating. However, addition

of silicon had detrimental effect on tribolayer formation and function during the sliding

process. It seems that Si somehow limits the formation of thick well-distributed tribolayer.

The higher wear rate and friction coefficient of Ti0.75Si0.11V0.03N in relation to

Ti0.65Si0.11V0.15N coating was in good accordance to the lower amount of “vermicular” debris

detected in the wear tracks. Although Ti0.82V0.15N and Ti0.65Si0.11V0.15N showed similar V

contents, the higher amount of transfer layer was formed in the case of Ti0.82V0.15N. The role

of Si is not clear, although it is possible that Si-O top layer hinders ion diffusion [15] and thus

protects the coating from oxidation required to form sufficient amount of lubricious vanadium

oxide.

The worn surface of Ti0.52Si0.20V0.14N film was found to be covered by fish-scale-like

agglomerates along the worn surface indicating severe shear deformation of the tribofilm

during the sliding test [8]. EDS analysis suggested the presence of V-O, Si-O and Ti-O oxides

in this zone. Additionally, EDS analysis showed significantly stronger peak of Si in relation to

the one found for the other TiSiVN films. Therefore, the oxide layer is dominated by Si-O.

Due to low fracture toughness large wear debris particles are detached during the running-in

stage; these particles are then either removed from the sliding contact (they were found on the

side of the wear track) or oxidized forming adhered oxide layer both on coating and ball

surface. However, due to the strong shear force that the ball imposes to the coating, and the

more brittle nature of Si-O-rich oxide layer, a stick-slip adhesion process will take place,

deforming and breakage the wear debris and giving rise to the above referred fish-like aspect.

Therefore, the tribological efficiency will be lower in this coating in relation to the other V-

containing films due to the combined effect of: i) lower fracture toughness, ii) less lubricious

V-O phase formed due the high Si content and iii) stick-slip adhesion wear mechanism.

Annex I


Figure 6 – SEM pictures of the wear track of: a) Ti0.65Si0.11V0.15N coating tested against Al2O3 ball, b)

Ti0.52Si0.20V0.14N coating tested against Al2O3 ball, c) and d) magnification of the wear tracks of Ti0.65Si0.11V0.15N

and Ti0.52Si0.20V0.14N coatings, respectively.

3.2.4 Wear mechanisms: Sliding against HSS balls

Low hardness of HSS balls compared to Al2O3 resulted in high ball wear and wide

wear tracks (see Figure 4 b)). In fact, one might expect that hard coatings will resist much

more to sliding against relatively soft steel ball, whereas Al2O3 counterpart should induce

severe wear damage. However, Figure 2 and Figure 4 clearly indicate opposite behavior, i.e.

higher wear with HSS balls.

Reference TiN and Ti0.80Si0.15N coatings exhibited slightly different wear track

appearance. Figure 7 shows the typical morphology of the worn surfaces of these coatings and

the corresponding wear scars of the balls. EDS analysis performed on the transfer layer of

TiN coating, which was preferential formed in the central part of the wear track, evidenced

material transfer from the ball to the coating surface. The transfer layer thus consisted of

titanium and iron oxides. However, no oxygen was detected by EDS outside of the center part

of unworn coating. Although the wear scars of the counterparts were covered with as-

Vermicular black

wear debris

Fish-like

wear debris

200 µm

5 µm

200 µm

20 µm

a) b)

c) d)

Annex I


deposited and oxidized material from the coating for both TiN and Ti0.80Si0.15N coatings, the

latter exhibited higher material transfer to the HSS ball and, particularly, larger wear debris

particles. Hard non-oxidized particles embedded into compact adhered tribolayer act as sharp

cutting edges causing abrasive damage and accelerating wear. On the other hand, relatively

thick and well-adhered layer on the ball wear scar protected ball material from further

damage; thicker layer built during sliding against on Ti0.80Si0.15N resulted in lower ball wear

(Figure 2 b)). It is clear here that higher amount of non-oxidized debris adhered on the ball

protects the ball itself but produces higher wear of the coating; the absence of well-developed

oxide layer on the coating surface is detrimental decreasing coating wear resistance.

Figure 7 – Wear tracks induced by HSS ball sliding in: a) TiN coating, b) Ti0.80Si0.15N coating. Wear scar on

HSS balls tested against: c) TiN coating, d) Ti0.80Si0.15N coating.

Absence of continuous tribolayer in the wear track was observed as well for the

coatings containing vanadium. Typical worn surfaces of the coatings (Ti0.82V0.15N and

Ti0.65Si0.11V0.15N) shown in Figure 8 were covered by stripes of adhered material, which

covered only small part of the wear tracks, see Figures 8 a) and b), indicating that the wear is

driven by a combination of adhesion and polishing wear. EDS analysis revealed that the

stripes were composed of fully oxidized (i.e. no nitrogen observed) coating elements mixed

Transfer

layer

200 µm 200 µm

400 µm 400 µm

a) b)

c) d)

Annex I


with iron originating from the ball. Outside of the stripes no oxygen was identified by EDS.

The phenomenon of adhesive wear behaviour was believed to be a result of hard film and

relatively soft bearing ball [16]. Nevertheless, some of the V-O oxide present on the stripe

zones should still provide a lubricious effect since a decrease of the wear rate with V

incorporation is observed. This behaviour is enhanced in case of Ti0.52Si0.20V0.14N coating

where an even higher wear rate is observed; its lower fracture toughness, due to the higher

silicon content, constrains the V-O phase formation decreasing even more its content on the

wear track. Ti0.82V0.15N film displayed the lowest wear rate of all the coatings due to the

higher amount of V-O oxide formed at the surface.

In summary, we demonstrated that V additions successfully reduced the wear rate and

friction coefficient of TiSiN system; moreover we also showed that the tribological behaviour

of films is strongly dependent on the counterpart material due to the presence of dissimilar

wear mechanisms.

Figure 8 - SEM pictures of the wear track tested against HSS balls of: a) Ti0.82V0.15N coating, b) Ti0.65Si0.11V0.15N

coating. c) and d) magnification of the wear tracks of Ti0.82V0.15N and Ti0.65Si0.11V0.15N coatings, respectively.

200 µm 200 µm

20 µm 20 µm

a) b)

c) d)

Annex I


4. Conclusion

The results showed that alloying TiSiN system with increasing V content improves the

resistance to wear and decreases the friction coefficient; however, the tribological behaviour

is strongly dependent of the counterbody ball material. SEM and Raman analysis revealed

that the influence of V was due to the formation of V-O lubricious oxide on the contact zone.

When Al2O3 balls were used in the tests, the wear was driven by polishing wear and oxides

containing V-O are spread all over the wear track. On the other hand, when tested against

HSS balls, adhesion wear took place with material transfer from the ball to specific zones of

the wear track, being the lubricious effect of V-O significantly reduced. Ti0.80Si0.15N film

displayed the lowest wear resistance among all the tested coatings due to its lower fracture

toughness which promotes the formation of bigger wear debris acting as strong abrasion

players.

Acknowledgments



para a Ciência e a Tecnologia –, under the project: PTDC/EME-TME/122116/2010, as well as

the grant (SFRH/BD/68740/2010).

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Annex I


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